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Full text of "Chemically derived ceramic composites"

;hemically derived ceramic composites 
By 

BURTRAND INSUNG LEE 



DISSERTATION PRESENTED TO THE GRADUATE SCHOOL 
OF THE UNIVERSITY OF FLORIDA IN 
PARTIAL FULFILLMENT OF THE REQUIREMENTS 
FOR THE DEGREE OF DOCTOR OF PHILOSOPHY 



UNIVERSITY OF FLORIDA 
1986 



ACKNOWLEDGEMENTS 

It is difficult for me to acknowledge everyone who yave helping 
hands during the tenure of my research as a graduate student at the 
University of Florida. 

It is hardly necessary to mention my advisor Professor Larry L. 
Hench for his inspiring support and stimulating thoughts. His sense of 
humor maintained a relaxed research atmosphere. He helped me to mature 
as an independent scientist. I am especially grateful for the freedom 
and independence he has given me throughout this research. li was in- 
deed a privilege and pleasure to work with him. 

Dr. Robin Sinclair at 3H Co. not only provided the facility to 
synthesize some of polysilanes but also taught me the techniques in 
polymer chemistry. The analytical services of 3M Co. which provided 
analytical data of the polysilanes are acknowledged with gratitude. Dr. 
Curt Schilling at Union Carbide is also acknowledged for his support by 
providing some of their experimental polysilanes. 

My thanks are extended to D. Dunnagan and U. Folz for thermal anal- 
yses, E. Jenkins for SEM, W. Acree for XRD, Dr. J. Newkirk and C. Turner 
for TEM, Dr. A. Gupta for GC, S. Yoon and S. Kong for proton NMR, Lester 
at the Engineering Machine Shop of the University of Florida, G. LaTorre 
for FT-IR, S. Kang for density and microhardness measurements, and 
Professor Batich and Dr. S. Kurinec for help in XPS. 

The financial support of the U.S. Air Force Office of Scientific 
Research through contract no. F49620-83C 0072 was certainly an essential 
part of my life at the University of Florida and is especially acknowl- 
edged as is the encouragement of Dr. D. R. Ulrich, contract monitor. 

ii 



I am also grateful to the members of my supervisory committee. 
Professors C. Baticn, M. 0. Sacks, L. Malvern, 0. Clark, and E. D. 
Whitney, for their advice and reading of the entire manuscript of this 
dissertation. 

On the nontechnical side, special personal thanks goes to my 
family. Instead of complaining about not spending much time with them, 
they rather were truly the Gatorade, "a thirst quencher." My mother-in- 
law, in particular, played too great a role to describe. 



m 



TA3LE OF CO.\T£NTS 

ACKNOWLEDGEMENTS i i 

LIST OF ABBREVIATIONS, ACRONYMS, INITIALISMS, AND SYMBOLS vi 

ABSTRACT ^ ^ 

CHAPTERS 

I. OVERVIEW OF CHEMICALLY DERIVED CERAMICS 1 

II. SILICON CARBIDE FROM 0R6AN0SILANE PRECURSORS 

Introduction 1q 

Experifnental 15 

Results 31 

Discussion 54 

Conclusions 96 

III. SILICON CARBIDE/SILICA COMPOSITES FROM CAR30SILANES 
AND ALKOXYSILANES 

Introduction luij 

Experimental Hj2 

Resul ts ] ■ 103 

Discussion 142 

Conclusions 152 

IV. SILICON CARBIDE/SILICA COMPOSITES FROM COMMERCIAL 
SILICON CARBIDE AND SILICON TETRALKOXI DE 

Introduction 155 

Experimental 153 

Results 167 

Discussion 213 

Conclusions 230 

V. OTHER CHEMICALLY DERIVED CERAMIC COMPOSITES 

Introduction 233 

Experinental 234 

Results 235 

Di s cu s s i on 238 

Conclusions 257 



TV 



Table of Contents (continued) 

Page 

VI. CONCLUSIONS AND RECOMMENDATIONS 259 

RE FE RE NCE S 264 

B lOGRAPH ICAL SKETCH 272 



LIST OF ABBREVIATIONS, ACRONYMS, INITIALISMS, AND SYMBOLS 

Ac Acetate ion or group 

AIBN Azobisisobutyronitrile 

A-PSS Allylic polysi 1 astyrene 

B.E. Binding energy in eV 

BPO Benzoyl peroxide 

-C=C or '^"y.v Vinyl group 

-C-C=C or ~\s Allyl group 

CDCl^ Deuterochloroform 

DgDg Deuterobenzene 

CFRI Chemical free radical initiator 

DCCA Drying control chemical additive 

DCP Dicumyl peroxide 

DMDCS Demethyldichlorosilane; Me2SiCl2 

DSC Differential scanning calorimetry 

DTGA Derivative thermogravimetric analysis 

EDS Energy dispersive x-ray spectroscopy 

en Ethyl enedi amine 

EtOH Ethanol 

FID Flame ionization detector 

FT-IR Fourier transform infrared 

GO Gas chromatography 

GPC Gel permeation chromatography 

IR Infrared 

J-PSS PSS prepared by Shinnisso Kako Co., Japan 

vi 



Me Methyl group, -CH3 

Hn Number average molecular weight 

MeCl2 Methylene chloride 

MPDCS Methylphenyl dicnl orosil ane, MeSiPhCl2 

Mrad Mega rad 

M.W. Molecular weight 

PC Polycarbosilane 

POMS Polydimethyl silane 

PDS Polydimethyl silane 

Ph Phenyl group, -C^jH^ 

PrOH Isopropyl alcohol 

PS Polysilane 

PSS Polysilastyrene 

PSS-0 Oligomer fraction of polysilastyrene 

P-PSS Petrarch's polysilastyrene 

R Reflectance, or diameter to length ratio of a fiber 

RT Room temperature 

T Temperature, or transmittance 

TEOS Tetraethylorthosil icate: tetraethoxysilane 

TMOS Tetramethoxysi lane 

TMS Tetramethylsilane 

SEM Scanning electron microscopy 

SS Silastyrene, oligomer fraction of polysilastyrene 

STEM Scanning transmission electron microscopy 

TEM Transmission electron microscopy 



VI 1 



TGA Thermogravimetric analysis 

TMA Thermonechanical analysis 

UV Ultraviolet 

v/o Volume percent 

ViSO Vinyl ic silane oligomer 

ViSP Vinyl ic silane polymer 

w/o Weight percent 

w/v Weight percent volume 

XPS X-ray photoel ectron spectroscopy 

XRD X-ray diffraction 

Greek Symbols 

a Linear coefficient of thermal expansion 

V Uncharged high energy electromagnetic quanta 
p Density, g/cc 

\ Wavelength 

V Frequency 

V Wavenumber 

A Differential value 
a Stress, strength 

Subscripts 

X Crystalline 

g Glass 

c Carbide 



VI 1 1 



Abstract of Dissertation Presented to the Graduate School 

of the University of Florida in Partial Fulfillment of the 

Requirements for the degree of Doctor of Philosophy 



CHEMICALLY DERIVED CERAMIC COMPOSITES 

By 

BURTRAND INSUNG LEE 

May iy8b 

Chairman: Dr. Larry L. Hench 

Major Department: Materials Science and Enyineering 

Silicon carbide was made from various organosilane precursors by 
crosslinking and pyrolyzing then in an inert atmosphere. Crosslinking 
of these silane precursors was studied by various means. The most suc- 
cessful means of crosslinking was found to be via a chemical free radi- 
cal initiator, dicumyl peroxide. The mechanism of crosslinking for the 
precursors was determined. 

Pyrolyses of the silane precursors were carried out and increased 
ceramic yields after crosslinking were shown as compared with uncross- 
1 inked precursors. The ceramic yields determined by TGA ranged from 10- 
70% depending on the precursors and the crosslinking treatments. 

Partially densified sol-gel derived silica monoliths were impreg- 
nated with the silane precursors while the silica monoliths were still 

ix 



highly porous. Diamond microhardness values increased 2-3 times from 
unimpregnated gel derived silica monolith. A noderace increase in frac- 
ture toughness, Kjq, and flexural strengths was achieved. Optical and 
mechanical properties, and porosity data are presented. 

Using sol-gel silica precursors and tne processing techniques of 
polysilanes to obtain 3-SiC, molecular composites of SiC witn Si02, with 
Ti02 and with AI2O3 were made in monoliths and powder forms. Monolithic 
composites with a molecularly dispersed SiC phase in the Si02 gel matrix 
showed a hardening effect by the SiC phase. 

The molecular composite powders of SiC/Al203 showed no or little 
crystallization of either phase after heating to 1400°C. 

Monolithic silicon carbide/silica composites were made using com- 
mercially available fibrous silicon carbides and a tetral koxysi 1 ane 
precursor. Modest to low flexural strengths were obtained after heat 
treating to 900-1400°C, because of the high porosity in the composite. 
Cold pressing of the SiC and silica sol slurry improved the density and 
flexural strengths. Notched 3-point fracture toughness values, Kt,-, was 
as high as 7 MPa^m-'-'^. Excellent thermal shock resistance and oxidation 
resistance of these composites are shown. 



CHAPTER I 
OVERVIEW OF CHEMICALLY DERIVED CERAMICS 



Ceramic materials are of critical importance in high technology 
(high-tech) areas where unique combination of properties, such as high 
strength, strength retention at high temperature, low thermal and elec- 
trical conductivity, high hardness and wear resistance, and high chemi- 
cal stability are required. However, because of their brittle nature, 
ceramic materials produce problems in design reliability in high perfor- 
mance structural applications resulting in catastrophic failures under 
stress. This inherent problem combined with poor cost effectiveness in 
fabrication of complex shapes severely limits wider applicability of 
current ceramic materials. 

Ceramic materials derived from chemical reagents have the potential 
to overcome these problems by 

1) Low temperature processing compared with traditional ceramic 
processing, 

2) Starting chemical compounds that can easily be purified to in- 
crease the purity of the ceramic materials, 

3) Having versatility in forming complex shapes and precise control 
of each step in the processing, 

4) Rendering homogeneous mixing; uniformity and, thus, reliability 
of the formed bodies can be improved, and 

5) A unique combination of microstructure and phase assemblages not 
obtainable by traditional ceramic processes may be obtained. 



Ceramic materials obtained by chemical processing are a rather 
recent development, despite the fact that the science behind the pro- 
cessing existed long before the ceramic applications were realized. 
Traditional ceramic science has been based more on physics and has been 
developed by optimizing the physical behavior with the Microstruct'jre of 
the material s. 

The term "ultrastructure processing" of ceramics has been intro- 
duced^ to represent the chemical manipulation and control of surfaces 
and interfaces during the earliest stages of formation in atomic or 
molecular scales. The so called "high-tech" ceramics are largely based 
on "ultrastructure processing" as this is one way in which engineering 
ceramics can potentially yield properties approaching the theoretical 
1 imit. 

Organic chemistry, once an anathema to ceramists, is recognized as 
a major source of new ceramic materials. By using ordinary chemicals as 
precursors to ceramic materials, one can study and control the chemical 
process in every step during the evolution of ceramics, from the start- 
ing chemical to the final product. Greater versatility in fabrication 
with more precise control of the process leading to dn extremely homo- 
geneous composite with superior properties is the goal and advantage of 
this approach. 

Ceramic fibers, optical glasses, ultrapure and ultrafine powders, 
and ceramic monolithic parts are some of the demonstrated materials 
derived chemical ly.i »2 Figure I-l summarizes some of the chemical pro- 
cesses. 



Add to powder 
as binder 



Carbo si lone 

Metal organ<c 

Metal olkoxide 

Metal salts 



Composite 

ceramtc 

body 



PyrolySiS or 

theffTKil 

decomposition 



Hydrolysis 
Polycondensalion 




Impregnate 
Porous ceramic 



cootinq/ttiin film 

Fiber 

Powder 



Impregnate 
Porous ceramic 
body 



— ^/ 
Partiotly- polymerized ' 
sol \ 



Use OS 
binder 



Metal carbides 
Metal nitrides 



Thermal 
Treatment 



Composite 

ceramic 

body 



Fibers 



_L 



Muiticomponent 

homogeneous 

noncrystalline gel 



Thermal 
Treatment 



Monolittiic porous 

structural 

Of reoctive 

gel powder 



sintered 
body 



Sintering 



Glass or 
Gloss ceramics 



Fig, I-l . Flow Diagram of Some of Chemically Derived Ceramics 
and Composites 



The sol-gel method, as shown in Fiy. I-l, is a notable example of 
obtaining oxide ceramics from metal -organic precursors. Pure monolithic 
parts, thin coatings, matrices for reinforced composites, etc. have been 
produced with controlled properties. Uranium oxide fuels were fabri- 
cated by the sol -gel method at Oak Ridge National Laboratory in the 
1970's.3 Active research is underway to understand the fundamental 
chemistry in the reaction steps, as well as in the appl ications. 2'"+ All 
facets of chemistry are involved. For example, a nuclear magnetic 
resonance technique has been found helpful in understanding the reaction 
mechanism of the sol-gel process. 2 

It has been shown2 that certain chemical additives change the phys- 
ical-chemical state during sol -gel transformation. The mechanism of how 
these additives function chemically is not fully understood. Various 
dopants may be added to the sol as a chemical reagent by forming a mo- 
lecularly homogeneous solution. 

Tne most understood sol-gel process is in the production of silica 
glass. Silica sol-gel reactions involve hydrolysis and polycondensation 
steps of a metal-organic precursor, as shown in eqs. I-l and l-Z. 

($4ijOR)4 + n HgO -.- = Si-OH + 4 ROH 
2 = Si-OH > = Si-O-Si = + H2O (1-2) 

The hydrolysis and polycondensation reactions initiate at numerous 
sites within the Si(0R)4 precursor + H2O solution as mixing occurs. 
They eventually form a three dimensional linkage of Si-O-Si in a sub- 
micron scale and are called sol particles. The sol particles come in 
contact to form a gel network. As the gel network is aged at an ele- 
vated temperature, the monolithic body is strengthened and becomes more 
like a ceramic body. 



Some other oxide ceramic materials derived chemically may be given 
below, 1) alumina from A1(0R)3 by Yoldas,5'5 2) lead titanate tiy Gurko- 
vich and Blum,'' 3) inaium tin oxide films by Arfsten et ai.,8 4) mono- 
sized SiOo and TiOo powders by Barringer et al.,^ and 5) single and mix- 
ed phase oxide powders by Mazdiyasni . 1° 

Similar to the sol-gel process of obtaining metal oxides, pyrolysis 
of organometal 1 ic precursors results in nonoxide ceramic materials of 
the constituent elements. Thus far, successful examples are silicon 
carbide (SiC) and silicon nitride (Si3N4) from polymers containing sili- 
con-carbon and silicon-nitrogen bonds in the backbone. "* Boron nitride 
and boron carbide can also be made from organometal 1 ic precursors.'* 
Titanium carbide, titanium nitride, titanium boride, silicon boride, and 
aluminum nitride may be possible from organometal 1 ic precursors.'* Table 
I-l lists ceramic materials that can be made from chemical processing of 
organometal 1 ic compounds.'* 

It is difficult to make complex shapes of dense refractory ceramics 
such as SiC or Si^^f^ using convem:ional high temperature sintering, hot 
pressing, or hot isostatic pressing methods without a sintering aid. 
Grain boundary phases are often introduced in materials, degrading high 
temperature performance and oxidation resistance. Refractory carbide 
and nitride fibers are nearly impossible to make using traditional pro- 
cessing methods. 

In making fibers from pyrolysis of an organometal 1 ic precursor, 
densification accompanies pyrolysis, which eliminates a separate sinter- 
ing process. By analogy to a carbon fiber made from a carbon polymer, 



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organometal 1 ic precursors offer a means to make refractory carbides and 
nitrides at potentially lower temperatures with the easy forming opera- 
tions of traditional polymers. 

One of the primary thrusts in applications of these materials is in 
gas turbine engines because these materials are strong and stable at 
temperatures no metal can withstand, and they also have high thermal 
shock resistance, which is necessary for heat engine components. Cer- 
amic materials derived from polymer precursors in the form of foam raay 
also be used for thermal insulation, filtration, and packing. Thin 
films may be applied in electronic devices and metal -ceramic joints. 
Extremely homogeneously doped high temperature semiconductors may be 
made this way as well. Boron nitride fiber made from organometal 1 ic 
precursors can be used as a dielectric material where alumina and silica 
fibers are less desirable. These organometal 1 ic precursor materials can 
also be used as binders in powder forming processes. There is also a 
high probability of obtaining infrared transmitting films of sulfides 
and selenides, superconductive fibers of NbN, NbC, silicides, sulfides, 
and borides via polymer precursor pyrolysis. 

Intrinsic flaw sensitivity and brittleness continue to impede the 
broader applications of monolithic ceramic components. These instrinsic 
weaknesses can be overcome by incorporating a high modulus, small diam- 
eter ceramic reinforcing phase in a ceramic matrix to change the failure 
mechanisms to tough, noncatastrophic modes. 

Composite materials on the ultrastructural level can be achieved by 
mixing polymer precursors containing constituent elements, e.g., a 



polysilane mixed with polyphenyl borazol e yields a composite of SiC/3N.3 
Polymers containing Si, C, and N can be used to obtain a SiC/Si3N4 com- 
posite.'* 

A reaction of polycarbosi 1 ane and Ti(0R)4 can yield a composite of 
SiC/TiC.i2 In this composite process Ti(0R)4 not only provides the TiC 
phase, but also crosslinks the polycarbosi 1 ane, hence maintaining the 
shape of the green body during the subsequent heat treatments and in- 
creasing the ceramic yield. The structural scale of these precursor 
based composites is in the 1-10 nm range, as compared with the 1 to 100 
]im or larger range of composites made by traditional processes. ^3 

Composites are leading a new era in structural engineering. The 
development of high performance materials and advances in fabrication 
technology are laying the groundwork for revolutionary changes in 
structural design. In order to go forward with high speed in ceramic 
composite technology, it is necessary for engineers to break away from 
engineering thought processes that have been developed over decades of 
working with conventional materials. 

Active research on ceramic matrix composites began no earlier than 
1982, according to Persh.^"* Even then, the ceramic matrix composites 
were based more on applied physics using the traditional processing 
methods, such as hot pressing matrix phase powder with a reinforcing 
phase. 

The objectives of this dissertation are thus based on an explora- 
tory study and development of new methods to obtain ceramic materials 
derived by chemical means. Processing and properties of SiC/SiOn 



composites utilizing techniques of sol-gel derived Si02 and SiC via 
organosil anes are the main topics of this work. Other ceramic compos- 
ites from chemical origins are also part of this dissertation. The pri- 
mary motivation and objectives for the work presented in this disserta- 
tion are an interest in the development of new processing methods based 
on chemical processes and an understanding of these processes. For cer- 
amic and composite materials in this work, the emphasis is more on con- 
cepts rather than the final products with exciting quantitative data in 
part because concepts are felt to be of greatest use to those developing 
ceramic composites. 

The more elaborate and topical introductions are given in the 
beginning of each chapter. 



CHAPTER II 
SILICON CARBIDE FROM ORGANOSILANE PRECURSORS 



Introduction 






Many ceramic materials have specific properties that make them 
ideal for energy related systems. Silicon carbide (SiC) is one of the 
leading candidates for high temperature structural applications because 
of its low density, high-temperature strength, chemical stability, re- 
fractoriness, high thermal shock resistance, and creep resistance. To 
achieve these desirable properties of silicon carbide, it is necessary 
to develop a reproducible and reliable method for producing the material 
in complex shapes and with a controlled ul trastructure. 

In the conventional process for producing SiC material, silica in 
the form of sand and carbon in the form of a coke are reacted together 
at 2400°C in an electric furnace. The SiC produced is in relatively 
large grains which are subsequently ground to the desired size.^^ 

The Cutler process^^ was developed to produce SiC material with 
superior properties and cost effectiveness by using rice hulls. From 
this process, the commercially known a-SiC whisker Silar" by AkCO is 
obtained. The major advantage of it is that it has a much lower pro- 
cessing temperature, ~1500°C, than the more traditional process. 

The increasing search for new types of high-strength materials, and 
for performance improvement in the existing ceramics, has pushed several 
nonconventional approaches to ceramic synthesis. 

10 



11 



As presented in Chapter I, obtaining nonoxide ceramics via pyroly- 
sis of organometal 1 ic precursors has potential advantages over the con- 
ventional methods of producing materials in low temperature process, 
higher purity, fabrication of complex shapes, greater homogeneity, new 
fabrication procedures leading to continuous fiber, coatings, and 
impregnated porous structures. At a more fundamental level, polymer 
routes can allow control over the microstructure of the intended ceramic 
product with a unique combination of microstructure and phase assem- 
blages and important consequences for both physical and chemical proper- 
ties.'* Some of the more important applications of nonoxide ceramic 
materials obtained via polymer routes are listed in Table II-l. 

The first use of organic polymers to produce an inorganic refrac- 
tory material was probably the development of graphite fiber from poly- 
acrylonitri le in late 1950. ^'^ Other ceramic materials from organometal- 
lic polymers v/ere first noted by Chantrell and Popper. 19 A partial 
history of the development of nonoxide ceramics from organometal 1 ic pre- 
cursors is given in Table II-2.'* 

However, early workers50"52 ■\q organosi 1 anes (OS) genuinely be- 
lieved that polysilanes were worthless and regarded them as undesirable 
by-products of a faulty synthesis. Tnis all changed in 197b when Yajima 
and his coworkers22-30 ,53-55 demonstrated that the polydimethyl si 1 ane 
that was regarded as an undesirable by-product by previous investi- 
gators50"52 is indeed a precursor to g-SiC. The reaction scheme, as 
shown in Fig. II-l, is the polymerization of dimethyldichlorosil ane 
[(CH3)2SiCl2] by dechlorination to yield polydimethyl silane (POMS). The 



12 



Table II-l. 

Some of the Demonstrated Applications of 

Nonoxide Cerauiics Derived fron Organonetal 1 ic Precursors 



Form 



Appl ications 



Fiber 



Reinforcement for composites 
weaves, wovens, mattes 



Mono! ith 
Foam 



Monolithic bodies for high temperature parts 



Filters, packing, insulation, heat exchanger 



Powder 



Press to bulk body, 
Fil ler material 



Thin Film 



High temperature electronic devices 



Thermosetting polymer Metal -ceramic, ceramic-ceramic joints, bind- 
er in powder forming 



13 



Table II-2 
A Partial History of Monoxide Ceramics Via Polymer Psoutes 







Precursor 


Ceramic 






Year 




Polymers 


Products 


Investigators 


Ref. 


1960 




phosphonitric 
chlorides 


P-N 


Ainger, Herbert 


18 


1965 




unknown 


BN, AIM, 
Si3N4, SiC 


Chantrel 1 , 
Popper 


19 


1974- 


■75 


polysilanes 


Si-C-N 


Verbeek and 

Winter et al . 


20, 21 


1976- 


■81 


polycarbosilanes 


SiC 


Yaj ima et al . 


22-3U 


1976 




polyphenylborazole 


BN 


Taniguchi , Harada 
Maeda 


31 


1978 




carboranesiloxane 


SiC-64C 


Rice et al . 


32, 33 


1979- 


■80 


polycarbosi lanes 


SiC 


Scni 1 1 ing, 
Williams, Wesson 


34-39 


1980- 


■81 


polysilastyrene 


SiC 


West et al. 


40 






polycarbosi lanes 


SiC, Si-C-N 


Baney and Gaul 


41-46 


1981 




polytitanocarbosilane 


Si-Ti-C 


Yaj ima et al . 


12 


1982 




polysilazanes 


Si-C-N 


Penn et al . 


47 


1982 




polysilazanes 


Si 3^4 


Seyferth, Wiseman 


48, 49 


1984 




vinylic polysilane 


SiC 


Schilling and 
Wil 1 iams 


39 



14 



c _ 



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a 
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c 
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polymer chain then is rearranged to make it more reactive for a thermo- 
setting condition. The rearranged polymer called polycarbosi 1 ane has 
alternating silicon and carbon atoms. The thermosetting or cross! inking 
carried out in air is a necessary step in the Yajima SiC fiber synthesis 
to maintain the fiber shape during the subsequent heat treatments. Dur- 
ing the heat treatment, hydrocarbon products are eliminated to yield a 
ceramic char of SiC. 

Since Yajima and coworkers22"30 »53-S5 developed a 6-SiC fiber with 
excellent mechanical properties from the polycarbosi 1 ane, there have 
been several other precursors potentially superior to the Yajima pro- 
cess. 3'+"'-+° These are vinylic silanes developed by Wesson and 
Wil 1 iams3'+"35 and Schilling et al.38.39 g.r\a polydimethyl phenylmethyl - 
silane, better known as polysi 1 astyrene (PSS) developed by West et al.'+o 

The vinylic silanes are reactive under a thermal crossl inking con- 
dition but, because of their low viscosity (liquid at room temperature), 
control of viscosity to draw fibers may require extra steps. On the 
other hand, PSS is a solid with good solubility in common solvents and 
possesses excellent tractability with good melt viscosity. However, it 
possesses no reactive functional groups for crossl inking. It was noted 
by West et al.'+o that the polymer strongly absorbs UV light at ~330 nm. 
Irradiation of UV with x~330 nm on PSS was shown to crosslink the poly- 
mer on the surface. '+0 However, the practicality of UV crossl inking of 
PSS for larger structural ceramics is in question. An alternate way to 
achieve bulk crosslinking is needed. 

It is important to note that vinylic silanes and PSS are poten- 
tially superior to polycarbosi lane because they can be processed without 



16 



the separate thermal rearrangement and oxygen crosslinking (Fig. II-l), 
as required in the Yajima process. Oxygen crosslinking undoubtedly 
introduces Si-O-Si in the network and ends up as silica in the final 
product s-SiC. In order for the potential advantages of vinyl ic silanes 
and PSS to be realized, it is essential that crosslinking methods be 
developed which will avoid oxygen in tne SiC lattice following pyrol- 
ysis. 

The term polymer used here is defined as any organic or organo- 
metallic compound that is not a monomer. However, in some specific 
cases, oligomers are distinguished from polymers. 

It is the objective of this work to investigate crosslinking 
methods and some applications of vinyl ic silanes and PSS precursors to 
SiC. In this chapter the following topics are investigated and dis- 
cussed: 1) synthesis and characterization of the polymers, 2) modifica- 
tion of the polymers for crosslinking, and 3) crosslinking and pyrolysis 
to obtain SiC. 

Experimental 
Preparation of equimol ar dimethyl phenylmethyl copolymer 

Reagent grade toluene for a solvent was dried with sodium metal in 
~1 g sodium per 1 1 toluene by refluxing for 24 hours followed by dis- 
tillation through a one way air sealed glass apparatus using mineral oil 
bubbler. 

A starting monomer dimethyldichlorosi lane {Me2SiCl2) from Aldrich 
Chemical Co. was purified by distillation using a trap-vacuum technique 
with liquid nitrogen. Methylphenyl dichlorosil ane (PhMeSiCl2) monomer 



17 



also from Aldrich Chemical Co. was vacuum distilled in a Yamato model 
rotary evaporator at 80°C. 

A reagent grade sodium metal bar was cut to 47.5 g and placed in a 
dry 2 liter 3-necked round bottom flask. These operations were carried 
out in a glove box with N2 atmosphere. The dried and distilled toluene 
(850 ml) was added to the flask and the polymerization reaction appar- 
atus was set up, as shown in Fig. II-2. 

Sodium and toluene were mixed by stirring and heating to obtain a 
molten mixture of sodium dispersed in the solvent. Heating was discon- 
tinued to add dichlorosi lane monomers. A premixed solution of 61 ml of 
MegSiCl?, 81 ml of PhMeSiCl2, and 50 ml of dry toluene was added slowly 
through the air sealed side arm, while stirring was continued and N2 gas 
was continuously flowing through the apparatus. 

The rate of addition of the premixed dichlorosil ane monomers was 
adjusted to maintain the gentle refluxing temperature of ~98°C, since 
the initial dechlorination reaction is highly exothermic. A typical 
duration of the monomer addition was ~30 min. Upon completion of the 
monomer addition, external heating was restored to achieve a gentle 
reflux. The reaction flask was kept dark by wrapping it with aluminum 
foil. The reflux continued for 1.5 hours before cooling the reaction 
mixture to room temperature and then poured into an isopropyl alcohol 
bath (PrOH) with stirring. 

Fractionation of the reaction products was carried out by first 
separating them in PrOH. The polymer fraction was precipitated out 
while the oligomer fraction remained in the solution. The oligomer 



18 



Silane 
mixture 



^2 out 




variable 
speed 
motor 



Thermometer 



ToIuene-»- Sodium 

+ Silane ^ v:.ir 



Heating mantle 



Fig. II-2. Apparatus for Polymerization Reactions for 
Synthesizing Polysilastyrene 



19 



fraction in PrOH was collected by distilling off the solvent by a rotary 
evaporator. The excess sodium residue was decomposed in PrUH. The oli- 
gomer fraction collected was redissolved in toluene and washed with dis- 
tilled water three times in a 200 ml separatory funnel to extract any 
residual salt product. Then the toluene solution of the oligomer was 
rotoevaporated to obtain a viscous oligomer fraction of PSS (PSS-0) 
which was kept in a brown bottle after vacuum drying and weighing. 

The PrOH insoluble fraction was washed with 200 ml PrOH and with 
200 ml EtOH twice after draining the PrOH by filtering. After the solid 
polymer fraction was dried in a vacuum oven for five hours at 55°C, it 
was redissolved in warm toluene and the toluene insoluble fraction was 
separated out. The toluene insoluble fraction was thought to be a 
highly crosslinked polydimethyl si lane. This fraction was washed with 
water five times in a separatory funnel and dried in a vacuum oven at 
80°C for ten hours. 

The toluene soluble fraction (PSS-P) in toluene solution was washed 
with water five times to insure that all unreacted Si-Cl is hydrolyzed 
out. This was done by titrating the effluent with AgN03 solution. Then 
PSS-P in toluene was reprecipitated in 7 liters of PrOH. The bright 
white precipitate was collected by filtration and dried in a vacuum oven 
at ~50°C for ten hours. A pure PSS-P should appear as a white powder or 
a clear, colorless solid. 

The variations in reaction conditions for the subsequent runs are 
given in Table II-3. 



20 



Table II-3. Summary of Reaction Conditions for 
Preparation of Polysi 1 astyrene and Product Designation, 





Mole Dichlorosilanes 








Product 


Mole Na 


Addition 


Reaction 


Reaction 


I.D. 


Volume Toluene 


Time, min. 


Time, hr. 


Temp., °C 



PSS-10 


0.5 mole each 


PSS-IP 


2.05 mole 




1 1 


PSS-20 


0.5 mole each 


PSS-2P 


2.09 mole 




1 1 


A-PSS-P 


0.5 mole MePhSiClg 


A-PSS-0 


0.4 mole Me2SiCl2 




0.1 mole Allyl MeSiCl2 




2.04 mole Na 




1 1 toluene 



4U 1.5 105 



30 2 108 



17 2 108 



21 



Infrared spectra for the polymerization products were taken by 
using a Perkin-E1mer IR Spectrophotometer Model 283 v^ith KBr Pellet in 
transmission mode and also by a Nicolet MX-1 FT-IR Spectrophotometer in 
diffuse reflectance mode. Proton NMR spectra were obtained by using a 
Varian XL-100 with CDCI3 or CgUg as solvents without the TMS reference. 
Molecular weight distributions of the products were determined by gel 
permeation chromatography (GPC) using polystyrene as a reference in THF 
solvent and using a refractive index detector. 

Other PSS samples were provided by Shinnisso Kako Co. of Japan 
through the 3M Co. (J-PSSl and J-PSS2), courtesy of Dr. R. Sinclair and 
also by Petrarch Systems, Inc. (P-PSS), courtesy of Dr. B. Arkles. 
Vinyl ic silanes were provided by Union Carbide, courtesy of Dr. C. 
Schilling. They are oligomer and polymer fractions of 



Me3Si-f Si 





■SiMe- 



Two kinds of siloxane substituted PDMS containing a hydride functional 
group or phenyl group were provided by Petrarch System, Inc., courtesy 
of Dr. B. Arkles. 

The structure of each polymer unit and physical state of all tne 
polysilanes used in this study are given in Table II-4. 
Crossl inking and pyrolysis 

For crosslinking via y-f'^y irradiation, PSS was melt-coated on thin 
stainless steel plates in glass test tubes with vacuum, Ar, M2, He, air. 



22 



Table II-4. Structure Forinulas of Polysilane Unit and 
Physical State of the Organosi 1 anes at Room Temperature 



Organosilane 



Structure 
Formul a 



Physical State 

At Room Temperature 



PSS-P 



-Me Me_ 



L-Ph Me-" 



White to dul 1 yellow 
sol id 



PSS-0 



,Me Me -, 
li —Si-- 



ll-l 



■-Ph Me-" 
(where m < n) 



Yel low to brownish 
viscous liquid 



Allylic PSS 
A-PSS-P 

A-PSS-0 



Me 

I 
Si -Si- 

I I 
Me Ph 



Light yel 1 ow sol id 

Brownish viscous 
1 i quid 



Vinyl ic Si lanes 
VI SP 

ViSO 



pMe^ 
MeoSi --Si 

LI. 



X L 



rMe-, 
Si 



SiMe3 Colorless clear 

viscous liquid 
Y 

Colorless clear low 
viscosity 1 iquid 



Siloxane PDMS 
Hydride 



-Si-O-Si- (SiMe2)8 



Yel lowish viscous 
clear 1 iquid 



Phenyl 



Ph 

-Si-O-ii — (Si 



(SiMe2)8 



Yel lowish vi scous 
cloudy 1 iquid 



23 



or NpO atmospheres. Sone portions of the silane precursors were dis- 
solved in benzene in glass test tubes and sealed for irradiation. The 
vinylic silanes were placed in evacuated Dorosilicate test tubes. The 
glass tubes containing silane samples were irradiated witn y-radiation 
from a *^'-'Co source at 1" distance for various lengths of time up to 29 
days at room temperature. 

Chemical free radical initiators (CFRI), benzoyl peroxide (BPO), 
aszobi si sobutyronitril e (AIBN), and dicumyl peroxide (DCP) obtained from 
Polyscience Co. were recrystal 1 ized from methanol before use. A few 
grams of silane were dissolved in 5-10 ml of benzene in a test tube or 
in a 3-neck round bottom flask and then the silane solution was degassed 
with an inert gas. After 30-60 min., a CFRI in the range of 3-10 wt% 
was added under an inert atmosphere and the crossl inking reaction was 
carried out with heating on a hot plate or by a heating mantle, as shown 
in Fig. II-3. The crosslinking reaction in the 3-neck flask was allowed 
to reflux for twelve hours with continuous stirring before cooling to 
room temperature. The crossl inked product was extracted and washed with 
methanol . 

Crosslinking via DCP was carried out also in a sealed Teflon con- 
tainer for ViSP, ViSO, and PSS. About 1-2 g of vinylic silanes were 
well mixed with 0.05-0.07 g DCP by a spatula under N2. Polysil astyrene 
was also mixed with DCP after the polymer was made into a thick solution 
in toluene. The silane + DCP mixtures were cured in an oven at 110- 
150°C after the containers were tightly sealed. Other portions of poly- 
sil anes were cured without DCP under the same condition. 



24 



N 



2 out 



Water 
^out 




Thermometer 



acting 
xture 



Heating mantle 
Magnetic stirrer 



Fig. II-3. Apparatus for Crossl inking of Polysilastyrene 



25 



For crossl inking via Pf^"^, a 1.2 x 10"* M solution of Pt^"^ was 
prepared by dissolving 3.1 mg of HoPtCl^'GHpO in 50 ml of a mixed sol- 
vent of acetone (20 ml), ethyl ether (16 ml), toluene (10 ml), and EtOH 
(4 ml). Crossl inking via Pt**"^ was tried for A-PSS, ViSP, and ViSO pre- 
cursors by adding 10"^-10"^2 [^gles of Pf*"^ to 0.05-0.4 g of the silanes. 

Crossl inkings of Petrarch's siloxane PDMS were carried out by 
adding drops of concentrated THF/H2O solution of Zn(Ac)2, SnCl4, 
[Co(en)3]2(S04)3, SnCl?, triethanol ami ne, ZrCl4, DCP, BPO, AIBN, Pt^"^, 
and ethanolic NaOH followed by heating and curing in an air sealed glass 
vial up to 15u°C or heating in N-^ gas up to 3U0°C. 

Detection and confirmation of crossl inking of the polymers were 
tested by using FT-IR, solubility in a solvent (benzene or THF), and 
fusion at ~200°C. Pyrolysis of the polymers was carried out in a high 
temperature furnace with an inert gas flowing at a rate of ~100 ml/min. 
and also in a DuPont TGA 951 Thermogravimetric Analyzer with N2 or Ar 
continuous flowing with a heating rate of lO'^C/min. Differential scan- 
ning calorimetry (OSC) using the same DuPont Model 951 was carried out 
in continuous Ar flowing with a heating rate of 5°C/min. 

Crosslinking reaction products of PSS via DCP were identified by 
GC's in order to elucidate the reaction mechanism. The product gasses 
were introduced into a Tracor GC 550 and an HP 5880A GC. The detailed 
experimental conditions and sample preparation are as follows: 

Approximately 0.1 g of PSS-IP was dissolved in 1.5 ml of degassed 
benzene, and then ~0.01 g DCP added and mixed in a 20 ml glass vial. The 
thick solution was vacuum dried to evaporate benzene at room temperature 



26 



for tv/o hours. The dried sample (PSS/OCP) was placed in a ylass vial 

with a rubber septun or in a Pyrex glass loop directly attacned to the 

Tracor GC injection port. The samples in the glass containers were 

heated to ~300°C by a Bunsen burner for 0.5-1 min. The product gasses 

were either directly introduced to the silica gel column of a Tracor 550 

through a valve or drawn by a syringe through the septum. The gas drawn 

by a syringe was dissolved in a methylene chloride (MeCl2) solvent; a 

few microliters of this solution was injected into the capillary column 

of an HP 5880A GC 

The GC parameters used dre given below. 

Instrument: Tracor 550 and HP 5880A interfaced to an HP 85 
computer 

Column: Silica gel 2.7 m 60-200 mesh and glass capillary 

Detector: Flame ionization 

Column T: 45°C and programmed from 80°C to 250°C at 20°/ 
mi n . 

Injector T: 180°C 

Detector T: 180°C 

Carrier Gas: N2 

The overall experimental conditions for crossl inking of various OS 

precursors are summarized in Table II-5. 

Infrared spectra, SEM micrographs, and EDS spectra were obtained by 

using a Nicolet MX-1 FT-IR Spectrophotometer and a JEOL model JSM-35C 

electron microscope, respectively. The assignments of IR bands are 

based on reference number 56 and are given in Table II-6. 



27 



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29 



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30 



Table II-6. Infrared Absorption dands of Polysilanes 



V. cm" Mode Shape, Intensity 



300 Si -Si bending ' weak 

420 H-Ph rocking weak 

400-480 Si-Si stretching 

700-800 Si-C stretching broad 

800-850 Si-CH3 rocking 

850-1000 Phenyl C-H bending 

1020 Si-CH2-Si wagging shoulder 

1100 Si-O-Si stretching sharp 

Si-Ph medium 

1250 Si-CH3 bending strong 

1400 -CH3 deformation broad, strong 

1480-1580 Aromatic C=C weak, sharp 

1620 OH bending of adsorbed water 

1600-1680 C=C aliphatic 

1710 C=0 medium 

1800 -Ph broad 

2100 Si-H sharp, medium 

2900 C-H stretching in Si-CH3 strong 

3050 C-H stretching in atomatic narrow 

3450 OH stretching broad 

3630 Si-OH stretching broad 



31 



Results 
Characterizations of the precursor silanes 

Some of typical IR, EDS, and NMR spectra are given in Fiys. II-4 
through 11-19. The M. W. distributions, X yield of each fraction, and 
MePhSi/Me2Si ratio for PSS are given in Table 1 1-7. The MePhSi/Me^Si 
ratios were estimated by integrating the deed under the peaks and nor- 
malized by the number of protons in each group of the peak. 

The oligomer fraction of PSS-1 in Fig. II-4 shows some C-OH and 
Si-OH (-3300 cm"^ and 3600 cm"M, Si-H (-2100 cm'M, possibly some un- 
saturated carbon, i.e. C=C (-1600-1900 cm"-^), strong and sharp Si-Me 
stretch (-1250 cm"-'-), the strong and broad band for Si-O-Si, and Si-Ph 
overlapped with Si-O-Si (-1100 cm"M, -Ph (-700 cm"^), and an Si-Si 
stretching band at -450 cm" . In PSS-IP, there is not as much Si-OH and 
little C-OH, less Si-H, and a small but sharp peak for Si-Ph at -1100 
cm"-'- is shown (Fig. II-5). 

Figure II-6 suggests that the oligomer fraction of PSS-1 has a more 
complex structure, indicated by multplets of the CH3 region (-1-2 ppm 5 
scale) and the Ph-H region (8-9 ppm), than the corresponding structure 
of the polymer fraction. It also shows a possible C=C band at -4.9 ppm 
which is absent in the polymer fraction (Fig. II-7). The peak at -2.6 
ppm may be ascribed to to -CH2- or to Si-H. 

Figure II-8 for PSS-20 shows that a larger proportion of Si-H is 
present in the oligomer, but less Si-OH and C-OH is present than in 
PSS-10 (Fig. II-6). Although PSS-10 and PSS-20 are both oligomer 
fractions of PSS, the IR spectra show that they are not exactly the same 



32 







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48 



Table II-7. M. W. Distribution, % Yield, 
and the Ratio of MePhSi/Me2Si for PSS 









% 


Yield 


Ratio 




M. W. 


(Mn) 


01 


iyomer 


MePhSi/Me2Si 




01 igomer 


Polymer 


01 igomer 


Silane 


Polymer 


Insol 


. Polymer 
72 


Polymer 


PSS-10 


390 




1.05 






103 




19 


1.1 


PSS-IP 


10 X 




8 




PSS-20 


344 


T 




72 

15 


1.20 

1.35 


PSS-2P 


11 X 


10^ 




11 




A-PSS-P 


331 


o 




75 
12 


1.1 
1.2 


A-PSS-0 


11 X 


103 




11 





J-PSSl 



6.8 X 10^ 



49 



compound; see qualitative differences at ~1720 cm" , ~1580 cm"-'-, ~1250 
cm"-^ , and -450 cm" . 

Figures II-9 and 11-10 show that the allylic PSS oligomer has a 
greater unsaturated carbon group, as indicated by the broader band at 
~16U0 cm"-'-; this is supported by the NMR data which show small humps 
around 3-7 ppm. However, this cannot be conclusive because the bending 
mode of water is also at ~1600 cm" . Tne band shape at ~145U cm"-*- is 
different from the nonallylic PSS. A significant amount of Si-O-Si may 
be present in A-PSS-0. Groups including possibly Si-H and an allyl 
group represent ~4% of the total protons based on the NMR data. Allylic 
PSS polymer represented by Figs. 11-11 and 11-12 contains a smaller pro- 
portion of Si-H relative to C=C. Carbon-13 NMR (Fig. 11-13) does not 
reveal any additional information. 

Figure 11-14 shows that only Si as a metallic element is present in 
PSS-IP under EDS analysis. 

Figure 11-15 for the solvent insoluble fraction of PSS-1 presumably 
due to the crossl inked network indicates that the polymer may be mainly 
composed of 

,Me Me 

i— a. 



Me Me^" 

An infrared spectrum of J-PSSl (Fig. 11-16) obtained from the diffuse 
reflectance mode of FT-IR indicates that Si-H is present, as well as 
-OH, but not as much Si-O-Si is shown. An IR spectrum for ViSO (Fig. 
11-17) obtained by the same way as for J-PSSl shows a larger proportion 
of Si-H. A small but sharp band at ~1600 cm~^ may be that of C=C. The 



50 



NMR spectrun of ViSO in Fig. 11-18 shows the presence of a vinyl group 
at 6.2 ppn. It is difficult to see the presence of Si-H, although snail 
humps between 1-4 ppm are shown. The NMR spectrun for ViSP (Fig. 11-19) 
indicates that the polymer has essentially the same structure as the 
ol igomer ViSO. 
Crossl inking and pyrolysis 

Si lanes exposed to a y-ray dose greater than 200 Mrad (~30 days at 
1" distance from the source at 0.3 Mrad/hr) were infusible at tempera- 
tures above 200°C and were insoluble in THF or in benzene, indicative of 
crossl inking. Additions of CFRI prior to y-ray irradiation made no 
difference in the crosslinking reaction rate. 

Fourier Transform IR spectra of PSS samples before and after 29 
days of y-ray irradiation are shown in Fig. 11-20. The polymer film on 
a stainless steel plate after 29 days of irradiation in vacuum showed 
insolubility in THF and infusibility upon heating up to 250°C. 

Gamma-ray irradiation of the vinyl silane oligomer in vacuum showed 
an increase in viscosity within twelve days from a watery fluid to a 
semisolid form. The vinyl silane upon heating at ~200°C for 10 min in 
N2 was transformed into a light yellow translucent solid which was in- 
soluble in toluene. 

Gamma-ray irradiation of PSS in a NoO atmosphere for >11 days 
changed the color of PSS from translucent yellow-green to bright red- 
brown. The viscosity of PSS decreased sharply, indicating the degrada- 
tion of the polymer. 

Among the CFRI's investigated (BPO, AIBN, and DCP), only DCP in the 
range of 2-15 wt% in PSS/benzene solution showed a positive crosslinking 



51 




2000 



1466 1199 932 
WAVENUMBERSlcm-') 



398 



Fiq. 11-20. Reflectance FT-IR Spectra of PSS Before and After 
29 Days of V-ray Irradiation in Vacuum 



52 



reaction. The CFRI DCP has the highest decomposition temperature of 
~150°C of all other CFRIs. The FT-IR spectrum of DCP crossl inked PSS at 
250°C compared with that of the as-synthesized PSS is shown in Fig. II- 
21. The DCP reacted polymers were insoluble in benzene and infusible at 
temperatures above 20U°C. The density of PSS crossl inked by DCP measur- 
ed by mercury volume displacement was 0.77 g/ml . Fourier Transform IK 
spectra of A-PSS samples after they were treated at different crossl ink- 
ing conditions are compared in Figs. 11-22 and 11-23. Oxygen cross- 
linking by heating in the air at 25-80°C came out negative for all PSS 
precursors. 

In the crossl inking of ViSP without DCP under the same conditions 
as with DCP, no solidification was observed within 20 hrs. However, 
curing at 30°C higher temperature 140°C resulted in solidification of 
the liquid ViSP, signifying crossl inking. 

The difference in the chemical structure of the ViSP samples cross- 
linked thermally compared with ViSP samples crosslinked with DCP is 
shown by FT-IR spectra in Fig. 11-24. 

Chromatograms of the gaseous products from a crossl inking reaction 
of PSS with DCP after being separated by a GC is given in Figs. II-2b 
and 11-26. Approximately 100 times more methane is produced as compared 
with ethane, as shown in Fig. 11-25. 

Differential scanning calorimetry thermograms are given in Figs. 
11-27 and 11-28 to compare the crosslinking mechanisms. In Fig. 11-29, 
DSCs of oligomer and polymer PSS are compared. Scanning electron micro- 
graphs of the surface of DCP crosslinked and pyrolyzed PSS and ViSP are 
shown in Figs. 11-30 and 11-31. 



53 




2000 1733 1466 1199 932 
WAVENUMBERS (cm-') 



665 398 



Fig. 11-21. FT-IR Spectra of PSS-1 and PSS-1 Crosslinked 
With 8% DCP 



54 




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CD 



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ViSP/DCP,110 C, 



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z 
< 

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UJ 

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u. 
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5600 4400 3200 2000 1400 

WAVENUMBERS (cm-*) 



800 



Fig. 11-24. FT-IR Spectra of Crosslinked Vinylic Silane 
Showing the Effect of DCP, Temp., and Time. 



57 











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tR, MIN 



B 



Fig. 11-26. Gas Chromatograms of the Reference and the Reaction 

Product From PSS/DCP Crossl inking Reaction Dissolved 

in Methylene Chloride (MeCl2). The Numbers on the 
Peaks Are the Retention Times in Minute. 



59 





V 

c 

Ui 



I 

O 

i-H 

o 




X 
Ul 



PSS- I P 

no OCP 




ex situ 



100 



200 



300 



400 



500 



600 



Temp. °C 



Fig. 11-27. DSC Thermograms of PSSl and PSS1/DCP Before and 
After the Crossl inking Reaction 



60 





■0 



I 





n 
u 




600 



Temp. ^ °C 

Fig. 11-28. DSC Thermograms of J-PSS2 and J-PSS2/DCP After the 
Crosl inking Reaction 



61 



a 

o 



r~> 



--^ 



PSS-IO/DCP 



J\ 



\ 



^ Ciosslinking , 

\ ■ 



I \ in situ 



y 



\ 



r 




100 



200 



300 



400 



500 



600 



Temp.^ °C 

Fig. 11-29. DSC Thermograms of PSS-1 Oligomer and PSS-1 Polymer 
Reacting in situ With DCP 



62 




Fig. 11-30. SEM Micrographs of PSS/DCP Showinx Blisters and Pores 

Generated by Gas Evolution. Top: After Crossl inking and 
Pyrolysis at 400°C in Vacuum, Bottom: After Crossl inking 
and Pyrolysis at 900°C in Nitrogen 



53 




Fig. 11-31. SEM Micrographs of ViSP/DCP Showing Pores and Surface 
Texture. Top: After ViSP Cross! inked With DCP at nO°C, 
Bottom: After Crossl inking and Pyrolysis at 900°C in N 



2 



64 



The more effective cross! inking conditions among the techniques 
tested on the various silanes are summarized in Table II-8. 

TGA thermograms of the precursor silanes are given in Figs. 11-32 
through 11-37 to show the char yield of SiC. Fourier Transform IR spec- 
tra of the pyrolyzed products are given in Figs. 11-38 through 11-42. 
In Fig. 11-39, the SiC product from PSS/DCP is compared with the commer- 
cial e-SiC Nicalon**. The spectra show the characteristic absorption 
band of Si-C stretching at 793 cm"! along with a small SiO^ band at 
-1040 cm"-^. An XRD powder pattern of PSS showed that the pyrolyzed 
product is amorphous which is identical with Nicalon®. 

The char yield of PSS without crossl inking was less than 20 wt%, 
which is close to the char yield of ViSO, while the char yield of the 
PSS/DCP systems show 52-61 wt% SiC. The char yields of pyrolyzed prod- 
ucts of SiC from various silane precursors are listed in Table II-9. An 
XPS spectrum for Si2P of SiC derived from ViSP in Fig. 11-43 shows -20 
atom% oxide silicon on the surface of SiC indicated by an overlapped 
peak at -107 eV B.E. 

Di scussion 



Based on IR and NMR data, PSS-IP contains a low level of Si-OH 
(-3400 cm"l) and Si-H (-2100 cm"^) bonds. Unsaturated carbon components 
(-1600 cm"-'-) may also be present at a low level. In PSS-10, significant 
amounts of Si-O-Si overlapped with Si-Ph, as shown by the broad absorp- 
tion band at -1100 cm" . The formation of Si-O-Si may be caused by 
water used to hydrolyze the residual Si-Cl in the polymer and the alco- 
hol solvent used for fractionation. 



65 



Table II-8. Summary of Effective Crosslinking Conditions 
Found for Different Si lane Systems 



Sil ane 



PSS-P 



Means 
Y-ray 

CFRI 



Effective Conditions 

vac. 29 days at 2.56 cm 
RT 

DCP 110-200°C 
in 10 min-10 hrs 
absence of oxyyen 



PSS-0 



CFRI 



DCP, 140-2bO°C 
in 2U min-12 hrs 
absence of oxygen 



A-PSS 



Thermal 

CFRI 
Pt4+ 



300°C in 20 min 
or 170°C in 14 hrs 
absence of oxygen 

IbO^C, 20 hrs 
absence of oxygen 

80°C, 12 hrs 



ViSP 



ViSO 



Thermal 

CFRI 

Thermal 
Y-ray 

CFRI 



150°C, 24 hrs 
absence of oxygen 

DCP, 120°C, 4 hrs 
absence of oxygen 

> 200°C, 72 hrs 

10 days, polymerization 
not a cross! inking 

DCP/110°C, ~3 hrs 
or 75°C, -12 hrs 
no oxyyen 



Ph 
SiOSiMeg 



CFRI 



DCP/N2, 300°C 
20 min 



66 




200 



400 600 

TEMP. °C 



800 



1000 



Fig. 11-32. TGA Char Yields of PSS-10, PSS-10 After DCP 
Cross! inking, and ViSO 



57 



100 1 



80 



UJ 

9 60 

CO 

111 
a: 



^40 

H 

X 

O 

LlI 

^ 20 




ex situ 
PSS/DCP 7% 

PSS 
ALLYL PSS 







_L 



200 



400 
TEMP 



600 



800 



1000 



Fig. 11-33. TGA Char Yields of SiC From P5S-1P, PSS-IP 
After DCP Crossl inking, and Allylic PSS 



68 




CD 

>- 

S- 
03 



3 
O 



00 



(J 

S_ 



X2 
3 



(U 

> 

o 



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3nais3y 



3,n 



69 



80 



§ 60 

CO 

u 

J ^0 1. 



20 - 











\ \ 




\ V 


- 


\ " " - _ _ ViSP/6% DCP 


^ 


^^ ViSO/6% DC? 






1 1 1 1 



200 



AOO 



600 
TEMP. 



800 



1000 



1200 



Fig. 11-35. TGA Char Yields of SiC for ViSP and ViSO After 
Cross! inked With DCP 



70 




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•r— 



Sil ane 



PSS-IP 



PSS-2P 



J-PSSl 
J-PSS2 

P-PSS 



A-PSS-P 



ViSO 



PSS-10 



77 

Table II-9. Char Yield of Pyrolyzed Products SiC 
froin Various Si lane Precursors 

Crossl inking method Pyrolysis Char Yield 

and Conditions Method % 

10% DCP in situ TGA 30 

10% DCP, 250°C TGA 61 
20 min, vacuum 

7% DCP, 250°C TGA 52 
10 min, N2 

5% DCP, 110°C TGA 23 
10 hrs, air sealed 

8% DCP, 250°C TGA 52 
20 min, N2 

- TGA 12 

800°C, 1 hr, N2 2 

5% DCP, 150°C, 12 hrs TGA 25 
air sealed 

5% DCP, 130°C, 12 hrs TGA 20 

air sealed 

TGA 15 

TGA 21 

Pt^^, 80°C, 12 hrs TGA 35 

5% DCP, 170°C 900°C, 1 hr. 22 

TGA 25 

4% DCP, 110°C, 6 hrs, TGA 25 
air sealed 

8% DCP, 150°C, 12 hrs, TGA 16 
air sealed 



78 



Table II-9 (continued). 



Sil ane 



PSS-20 



Crossl inking method 
and Conditions 

7% DCP, 300°C 
2U min, N2 



Pyrolysi s 
Method 

TGA 



Char Yield 

% 

67 



Insol . PSS 



TGA 



Insol. A-PSS 



TGA 



ViSP 



thermal , 150°C, 24 hrs 
3% DCP, 120°C, 3 hrs 
6% DCP, 130°C, 4 hrs 



920°C, 1 hr, iN^ 55 

920°C, 1 hr, N2 54 

TGA 72 



Si-0Si-Me2 



TGA 



<4 



Ph 

I 
Si-0Si-Me2 



8% DCP, 250°C 



TGA 



27 



79 




o 



o 

siNnoo 



m 
I 



80 



All of NMR spectra show slight excess of the PhSiMe unit in the PSS 
copolymer chain over the MeSiMe unit, despite the intention to form an 
equimolar copolymer. This must mean that the PhSiMe monomer unit is 
more reactive than the MeSiMe unit during the polymerization reaction. 
This means that in order to achieve an exactly equimolar copolymer of 
PSS, one would have to use a slight excess of Me^SiClj monomer. 

Hydrogen directly bonded to silicon in PSS comes from the fact that 
the polysilane chain ends are probably anionic in the sodium/toluene 
milieu and abstract hydrogen from the alcohol that was added to quench 
excess sodium. Since the Si-H is hydrolyzable to give Si-UH, when tne 
Si-H ends find Si-OH ends, they form a Si-O-Si linkage by a condensation 
reaction. „ 

A 

Although early workers of polysi 1 ane5°"52 considered that 4-Si4- from 

VJ '" 
Me 

MeSiCl2 was useless because of its insolubility, introduction of a 
phenyl group improves its solubility and tractabil ity. 

As given in Table II-6, all polymerization reactions carried out 
produced ~3 times larger oligomer fractions than polymer fractions. 
Although it is generally thought that polymer fractions are tne desired 
product in these reactions, tne oligomer was successfully repolymerized 
and crossl inked by using a DCP as a CFRI. This is an especially import- 
ant finding since the oligomer fraction in a liquid state at room temp- 
erature is more convenient to impregnate porous ceramic bodies in order 
to strengthen them. David, ^6 who pioneered PSS, has unsuccessfully 
tried to repolymerize the oligomer fraction by restarting the polymeriz- 
ation reaction with sodium in toluene, sodium chloride, lithium metal, 
lithium t-butoxide, potassium t-butoxide, or sodium with biphenyl as an 



81 



electron transfer agent. Ddvid56 clearly concluded that "all one can do 
is to fractionate them out of a polysi 1 astyrene product and discard 
them, . . ." Petrarch System, Inc., the only company that makes PSS 
commercially in the U.S., follows this practice. 57 

Oligomers of PSS are believed to be in cyclic form so that the 
crossl inking mechanism is expected to be different from the polymer 
fraction by opening up the six membered ring. This is partly evidenced 
by a higher TGA yield (67%) than that of a polymer (Figs. 11-32 and 
11-33). The DSC data (Fig. 11-29) show that oligomers require a higher 
temperature for crossl inking than polymers and the decomposition begins 
at ~50°C lower temperature than polymers. The high char yield of the 
oligomers, however, was partly caused by the loss of volatiles through 
the vacuum line during the ex situ crosslinking reaction with UCP. 

The longer reaction time of PSS polymerization, e.g. ID hrs, was 
believed to degrade the formed polymer by excess hot sodium. However, 
it appears that a longer reaction time than ~2 hours could have been 
used to increase the polymer fraction of PSS. It is not yet completely 
clear to what extent the longer reaction time improves the yield of the 
polymer fraction. A systematic study of the effect of reaction time on 
the M. W. of the silane products must be continued. 

In the allylic PSS polymer, small amounts of Si-O-Si and Si-H are 
shown (Fig. 11-11). The unsaturated carbon component, probably from the 
allyl group, is shown at ~1600 en"-'-. The mole ratio of MeSiPh to MeSiMe 
is 1.2:1. The i^C NMR (Fig. 11-13) did not reveal any more evidence of 
the presence of an allyl group in A-PSS-P. Some indications of allyl 
groups in A-PSS-0 are also shown in Fig. II-9. The group representing 



82 



Si-H and the ally! is shown as -4% of the total protons in A-PSS-P (Fig. 
11-12). Apparently, not all the allyl methyl dichl orosi 1 ane ended up in 
the products. A further study to account for this is needed. 

The low solubility of J-PSSl, J-PSS2, and P-PSS in the solvents 
must be a result of incomplete fractionation of the insoluble high M. W. 
fraction. This variation in the M. W. distribution in a polymer is most 
likely affected by the fractionation procedures. The consequence of 
this was observed in the different behavior of PSS under the same cross- 
linking conditions used in this study. 

The average number M. W. (Mn) for different PSS precursors are in 
close agreement, except that of J-PSSl. The GPC chromatograms show that 
the PSS polymer is bimodal with a ~4 times larger lower M. W. 
portion (Mn 9 x lO^) than the higher M. W. (Mn 3 xlO^). 

The very high M. W. fraction that is insoluble in common solvents 
and infusible upon heating, appears to be polydimethyl sil ane with some 
Si-H, Si-Ph, and Si-O-Si, as shown by the IR spectrum in Fig. 11-15. 
The 1 inear 

Me 




linkage apparently has been crossl inked via bridging oxygens of Si-O-Si 
type. The TGA char yield of this fraction was <10% (Fig. 11-34), which 
agrees with the result of Yajima et aU^"* 

In vinylic silanes, a large and sharp Si-H band at ~2080 cm"-^ is 
shown for ViSO in Fig. 11-16. The vinyl double bond appeared at ~1650 
cm"-^. However, the hydride proton is not apparent in NMR (Fig. 11-18). 



83 



This should be due to the relatively low concentration of the hydride, 
but this obviously disagrees with the IR data. This point will be 
discussed further in the latter part of this section. 

An energy equivalent to 200 Mrad or greater y-ray irradiation 
required to crosslink PSS is not unusual because of the phenyl group^s 
on the chain and the absence of an Si-H functional group. 

In Fig. 11-20 the effect of y-ray irradiation on the PSS structure 
after 29 days is shown and the sharp bands at ~70U cin"-^ representing the 
methyl group are lost. The DCP reacted PSS lost most of its IR bands 
for Si-CHj, and Si-H (Fig. 11-21). 

Tne enhanced crossl inking of carbon polymers in NpO under y-irradi- 
ation, as shown by 0kada,59 did not occur with silanes but rather the 
opposite was observed. 

Among the several CFRI studied, only DCP yielded an insoluble and 
infusible solid of polysilanes. This is probably due to the active 
methyl radical, which was not present in any other CFRI used. In the 
crossl inking reaction, DC? has to be decomposed to give the methyl 
radicals. This occurs at ~150-2G0°C. 

Allylic PSS is observed to be more reactive under crossl inking con- 
ditions than PSS. As shown in Fig. 11-22, A-PSS can be crossl inked 
thermally or via use of the CFRI DCP. For thermal crossl inking, a temp- 
erature >170°C is required for complete reaction. Allylic PSS is also 
shown to be crossl inkabl e by Pt'^''" catalyst. At least 2.4 x 10"^ mole 
Pt'^ per ~0.3 g A-PSS was required for an effective crossl inking of 
A-PSS, as shown by IR spectra (Fig. 11-23). 



84 



The crossl inKing reaction of A-PSS must be between Si-H and C = C, 

as shown in equation II-l. 

catalyst 
= SiH + C=C > = Si-C-C-H (II-l) 

The coupling reaction is catalyzed by Pt^"*".60 jhe TGA yield of A-PSiJ 
without a precrossl inking treatment (21%) is still greater than that of 
PSS-P (12%). After being crossl inked with Pt"^"^, the yield increased to 
35%. Differential scanning calorimetry (Fig. 11-44) shows that the 
crosslinking between Si-H and \= occurs at ~240°C. The CFRI DCP re- 
quires temperatures >170°C for complete crosslinking, but Pt"^"*" catalyzed 
the reaction at a temperature of ~80°C. This crosslinking reaction by 
Pt^"*" is unique to A-PSS and demonstrates the advantage of incorporating 
an allyl group into polysilane synthesis. 

Although a monomer with an Si-H functional group was not added in 
the A-PSS synthesis, the small amount of Si-H at the chain ends was 
still shown to be effective in the coupling reaction. However, inten- 
tional small amounts of a monomer with a Si-H functional group, e.g. 
H 

CI Si — Me should improve the crossl inkabil ity even further. 
2 

Vinylic silanes can be crosslinked both thermally and via the CFRI 
DCP. With DCP, the crosslinking is achieved faster and required a lower 
temperature: 110°C for ~4 hours as compared to 150°C for 12 hours with- 
out DCP. Without DCP, 150°C for 12 hours treatment still did not pro- 
duce complete crosslinking, as shown by the large Si-H IR peak in Fig. 
11-24. This is also shown in Fig. 11-45 with an expanded scale. 

The vinylic silanes (ViSP and ViSO) received from Union Carbide 
were reported^^ to have both functional groups Si-H and V\ 



85 





s 











500 



600 



Temp. °C 



Fig. 11-44. DSC Thermogram of A-PSS-P Showing that the Thermal 
Crosslinking Occurs at '-240°C 



86 



0001 




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30NV133nd3d % 



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87 



Although NMR confirmed the presence of ^ , the presence of Si-H is 
not certain. If ViSP and ViSO contain Si-H and ^ functional 
groups, they should be crossl inkaole via Pt^"*" even more readily than 
allylic PSS because \\ is more reactive than \_ . However, this 
was not observed. The vinylic silanes with a Pt"^"*" concentration greater 
than 9 x 10"'' mole per 1 g si lane and T > 8U°C for several hours did not 
solidify the liquid silanes. This is a question which cannot be answered 
unless further work is performed. 

The size of the Si-H peak in the IR spectra cannot directly be used 
to estimate the degree of crossl inking because of the high bond energy 
(-314 KJ/mole).5'+ This is shown by the Si-H IR band at -2080 cm"^ in 
SiC after pyrolysis at both 1000°C (Fig. 11-40), and 900°C (Fig. 11-41). 

Morterra and Low^z also observed the growth of the Si-H absorption 
peak when methoxylated aerosil was heated in a vacuum up to 750°C, while 
the absorption peak for the -CH3 stretching band at ~3000 cm"-*- decreased 
as the length of heat treatment at 750°C in vacuum increased. 

Nevertheless, it is shown in Figs. 11-22, 11-24, 11-45, and 11-46 
that the degree of crossl inking appears to be a function of the Si-H 
peak size. This is another area that needs to be further investigated. 

A difference between vinylic and allylic silanes under crosslinking 
conditions is in the reactivity of the functional groups. Vinyl groups 
are more reactive than allyl groups by vinyls forming more stable radi- 
cal intermediates. This was shown by the lower temperatures needed to 
crosslink vinylic silanes. This advantage is somewhat curtailed by a 
greater tendency of vinylic silane to be oxidized. Thus, one should 



88 



|PSS/28DAY 
VRAY 




Fig. 11-46, 



2600 2000 1700 ,1400 
WAVENUMBERS(cm') 

FT-IR Spectrum of a Region Showing the Fffect of 
Crosslinking on Si-H Band Intensities at -2080 cm" 



1 



89 



expect that in synthesizing more reactive silane precursors for cross- 
linking there is the danger of introducing none oxygen contaminant in 
the polymer and, thus, in the pyrolyzed product. 

In Fig. 11-46, the as-received PSS shows a sharp and strong absor- 
ption band for Si-H at ~2100 cm"-'-, a medium sized band for y-ray irradi- 
ated PSS, and a small band for 10 wt% DCP treated PSS. This means that 
PSS crossl inking can occur between Si-H's (Fig. 11-27), as well as Dy 
methyl free radicals from OCP at nigher temperatures. The bond energy 
of E C-CH^-H is 418.4 KJ/mole63 and the bond energy of a Si-CH2-H should 
be a little less than that of = C-CH2-H because of a greater electropos- 
itivity of a Si atom than that of C atom. Still, the bond energy of 
Si-H (314 KJ/mole)5'+ is much smaller than that of = Si-CHg-H, hence the 
crossl inking of PSS by DCP proceeds with Si-H bonds breaking at ~150°C 
followed by formation of Si-C-C-Si linkages via methyl radicals at 
~250°C. In Fig. 11-27, as-synthesized PSS shows a small exothermic peak 
at ~16U°C, which probably corresponds to the crossl inking reaction via 
Si-H. The DSC for J-PSS2 in Fig. 11-28 before the DCP treatment shows 
negligible crossl inking via Si-H coupling during the heating schedule of 
the DSC. Rearrangement of the polymer chain is thought to occur at 
~400°C. The small spikes at 100°C correspond to water evaporation. The 
reason for sharp endothermic peaks at ~520°C is not known, but it is 
thought to be the evaporation of a fraction of low volatility. After 
the endothermic rearrangement, decomposition to eliminate H2, CH4, CgH^, 
etc. actually begins to occur at ~420°C. 

The minimum at ~200°C for PSS with in situ DCP crossl inking (Fig. 
11-27) must be due to the decomposition of DCP. Under the heating rate 



90 



of DSC (5°C/min), the DCP crossl inking reaction may not be able to keep 
up with tne heating rate. The exothermic reaction was incomplete until 
the temperature was ~220°C. This supports the previous observation of 
incomplete crossl inking with DCP at temperatures below ~200°C and the 
low TGA char yield of the in situ DCP crossl inked PSS (Fig. 11-37). 

In the DCP precrossl inked J-PSS (Fig. 11-23) no chain rearrangement 
is evident. Instead of rearrangement, the decomposition begins at a 
slightly lower temperature, ~400°C. Tnis may mean that the molecular 
rearrangements have occurred during the preceding DC? crossl inking 
reaction. Tne primary chain rearrangement is probably a Kumada type, 6"+ 
as shown in equation II-2. 

Me H 

I / 

= Si-Si = > -> Si-CH2-Si = (11-2) 

A pyrolysis GC study of PSS (Figs. 11-25 and 11-26) with DCP showed 
a large amount of methane and acetophenone, which are some of the pro- 
ducts from the proposed crosslinking reaction given in Fig. 11-47. The 
amount of methane is too much to come from the Si-H coupling alone. 
Thus, the methane must be formed by the methyl radicals of DCP after ab- 
stracting methyl hydrogen from Si-CH2. The possibility of crossl inkage 
via Si-Ph-Ph-Si is doubtful because of the greater bond energy for-O^H 
(112 Kcal/mol) than for -CH2-H (104 Kcal/mol).55 

During crosslinking and pyrolysis, a precursor polymer is decom- 
posed and fragmented, preferably with free radical modes to achieve a 
high ceramic yield. However, pyrolysis conditions strongly affect the 
density of the pyrolyzed product. Density is increased during 



91 



Ph-C-00-C-ph ^-^A§-^ ph-C-0 



Me Me 



q 

-> Ph-f-ME + Me« 



riE 



Me Me 



f^E CH< 



Me» + 4Si-Si47. "^^S?^ 4^1— Si^ + CHi^ 



Me f^h 



>i Ar 



■Y J N 

Me Ph 



Me CH' 



Me Ph 



2e^i-Si^.. ^ ^i-Si-h. 



Me Ph 



Me CH- 



I[1e Ch2 

-fSi-Si-4^ 

k Ph 



Fig. 11-47. The Crossl inking Mechanism of PSS by DCP 



92 



pyrolysis but not usually to the tneoretical value due to the porosity 
generated by evolved yasses during tne process. Figures 11-30 and 11-31 
show the uneven surfaces and pores formed on PSS and ViSP after cross- 
linking and pyrolysis at 900°C in vacuum or in N2. The blisters formed 
on a sintered mass of SiC that are created during the crossl inking pro- 
cess are in subnicron sizes and remain after pyrolyzing up to 900°C. 

Although the SiC formed from organosi 1 anes and pyrolyzed at ~1000°C 
is known to be the beta phase, the x-ray diffraction pattern of the pow- 
der is amorphous because of its extremely fine grain size (~3 nm).2'+ 

Most TGA data for char yield are for samples after being ex situ 
precrossl inked in either N2 or in a vacuum. The ^1% char yield of PSS-0 
suggests that some of the most volatile fractions escape during the 
crosslinking reaction. However, the char yields of crossl inked PSS-0 
with respect to the PSS-0 starting weight are in the range of 35-45%, 
which is close to the theoretical yield of 45%, according to equation 
II-3.'*o 

Me Me 

-► Ph-H + CH4 + 2H2 + 2SiC (1 1-3) 






h Mr'" 



This demonstrates the potential usefulness of PSS oligomers as a filler 
phase for porous ceramic bodies by infiltrating the pores as small mole- 
cules followed by crosslinking and pyrolysis in the pores of the ceramic 
body. 

A higher char yield of A-PSS than that of PSS without precrossl ink- 
ing treatments is thought to be a result of crosslinking between Si-H 



93 



and S= during the pyrolysis, as discussed previously in this chapter. 
This in situ thermal crosslinking also applies to ViSO and ViSP. 

In the case of the more reactive ViSP, temperatures above 11U°C 
with UCP and above 150°C without DCP are required for crosslinkiny 
(Table 11-9). The Petrarch's H-Si-Si-0 substituted POMS was difficult 
to crosslink by all methods tried. This raises a doubt that the polymer 
has a Si-H functional group, or that a Si-U group may somehow effec- 
tively inhibit the functionality. No characterization to determine the 
structure was carried out on tne Si-0 linked polysilanes. Other evidence 
to increase doubt of the presence of a Si-H group is in the essentially 
nil TGA yield of the polymer upon pyrolysis. It may be the Si-0 which 
makes the crosslinking very difficult. 

The Petrarch's phenyl substituted Si-Si-0 POMS was crossl inked with 
difficulty by DCP which increased the TGA yield to 27%. The cross- 
linking mechanism of this polymer by DCP is presumed to be similar to 
the PSS/DCP system, but a Si-0 group must be suppressing the free radi- 
cal crosslinking. The usefulness of this polymer in ceramic applica- 
tions is presently unknown, i.e. the effect of Si-0 substitution needs 
to be investigated further. 

The TGA yield of in situ DCP crosslinked PSS-IP is -30%, which is 
lower than the theoretical yield of 45%. This low yield is due to the 
constant rapid heating rate (10°C/min) in the TGA operation so that the 
crosslinking temperature in the ~200°C region is passed in a few min- 
utes, which is not enough time for complete crosslinking. 

A pyrolysis study of PSS was carried out by Sinclair's at 400°C in 
vacuum. Each fraction he obtained is shown quantitatively below. This 



94 



distribution of PSS fractions suggests that the crossl inking of PSS via 
DCP allows recovery of fractions up to sone of median volatility (54.2%) 
as SiC product. The char yield of ~50% is in fact the inaximum theoreti- 
cal yield. However, without crossl inking, only the nonvolatile solid 
and maybe some of the low volatile solid fraction (~10%) is recovered as 
SiC upon pyrolysis. 



Fraction 

Highest volatility 

lost through vacuum line 

High volatile 1 iquid 

Medium vol atile 1 iquid 

Low volatile sol id 

Nonvol atile sol id 



Wt% 
12.4 

31.7 
17.7 " 
28.6 > 
7.y J 



54.2 



As the organosil anes were heated in a quartz tube furnace in N2, 
yellow gas was generated and condensed on the colder region of the tube 
wall. This same gas corroded the DSC metallic sample chamber, TGA 
sample boat made of platinum, Nicalon® continuous SiC fiber in SiC/SiC 
composites, and the thermocouple of a furnace. The yellow corrosive gas 
was first suspected as a chlorine compound, e.g. CI2, HCl , HOCl , etc., 
formed by the residual chlorine from the starting dichl orosil ane mono- 
mers. The residual chlorine was tested by dissolving PSS in benzene and 
hydrolyzing any Si-Cl with water by shaking the benzene phase and water 
phase in a separatory funnel, then the aqueous phase was titrated with 
AgN03 solution. No precipitate was observed. Based on this test, the 
possibility of the yellow gas as a chlorine compound was ruled out. 



95 



Silicon and/or SiC, at low oxygen partial pressure, can form sili- 
con monoxide at elevated temperatures. ^ 5»67 jt is then possible that 
the yellow gas is SiO, but the chemical reactivity of SiO with respect 
to metals, SiC, and SiOg is not known. If it were SiO, the low partial 
pressure oxygen source must be in the inert gas, e.g. a commercial re- 
search grade No gas contains ~1 ppm O2 and ~1 ppn H^O, and also in the 
organosilane itself. 

Figure 11-38 shows a large absorption peak for Si02 at ~1073 cm" . 
This large absorption band of SiOp is caused by a Si02 film on the sur- 
face of a monolithic PSS/SiC sample after crossl inking and pyrolysis at 
960°C. This surface oxidation must be caused by oxygen in the inert gas 
and/or during handling in air prior to pyrolysis. An FT-IR spectrum of 
the same PSS/SiC sample after the monolith was crushed into fine powder 
was taken again. This is shown in Fig. 11-39, along with commercial 
6-SiC Nicalon®, which is made from polycarbosil ane. The large Si02 
absorption band previously shown is diminished to a small hump at ~1U50 
cm"-'-, while the bands corresponding to Si-C at -800 cm"-'- remained the 
same. 

In Figs. 11-40 and 11-42 for J-PSS/SiC and ViSP/SiC pyrolyzed at 
900-1000°C, it can be seen that some type of organic residue still 
remains, as well as Si-H groups. This may be either from an incomplete 
transformation of the organosil anes to 6-SiC or from contaminants from 
external sources during handling prior to pyrolysis. This means that 
temperatures above 900°C are required to convert an organosilane com- 
pletely to B-SiC. However, an FT-IR spectrum of a ViSP/SiC monolith 



96 



pyrolyzed at 950°C and surface polished with SiC grit followed by wash- 
ing with acetone shows an essentially identical spectrun as Nicalon*® 
(Fig. 11-43). Therefore the organic residue nust be concentrated on the 
surface. 



Conclusions 

It has been shown that these various organosil anes as polymers and 
oligomers can be successfully converted to 8-SiC after crossl inking 
treatments and nearly reach the theoretical yield upon pyrolysis. The 
process of SiC production from silane monomers and PSS is summarized in 
Fig. 11-48. Tne applications of these materials and techniques for 
making ceramic composites are the subjects of the following chapters. 

The exploratory organosilane (OS) precursors used in this work have 
the potential to be formed into desired shapes using conventional low 
temperature plastic processing, then pyrolyzed to obtain SiC material 
with a yield that is nearly the theoretical limit. These precursors can 
also be used to impregnate porous ceramic bodies. 

The highest char yield is given by the polymer fraction of vinyl ic 
silane (72%). The allyl group on the polysilane chain improved cross- 
linkability. A complete crossl inking of OS precursors increased the 
char yield of SiC. For complete crossl inking, OCP with temperatures 
greater than 250°C for PSS and greater than 130°C for vinyl ic silanes 
are required. Other combinations of the functional groups such as Si-H, 
— ^ , — v^ , <\ , etc. should further improve crossl inkabil ity and 
ceramic yield. 



97 



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Char yield of SiC is roughly shown to be a function of concentra- 
tion of the crossl inking agent DCP, and a fjnction of temperature and 
time. However, temperature has a much greater effect on the degree of 
crossl inking than time. The possible oxygen contamination from the 
crossl inking agent, DCP was not observed. 

These OS precursors are prone to oxidation during handling, cross- 
linking and pyrolysis to include oxygen contaminant in the final SiC 
product. For a complete freedom from oxygen contamination, there must 
be more convenient ways to process the OS precursors, e.g. a glove box 
fully equipped to avoid oxygen. 

In order to increase the polymer fraction of PSS, a longer reaction 
time than two hours would be desired in the synthesis. 

Variables that may have significant effects on the PSS polymer 
yield and quality are 1) rate of chlorosilane monomer addition, 2) 
amount and kind of solvent relative to the reactant, 3) state of disper- 
sion and the amount of sodium, 4) reaction temperature, 5) reaction 
time, 6) reactant's molar ratio, and 7) environmental factors such as 
oxygen, ultraviolet light, etc. 

To improve the lot-to-lot reproducibility of the processing of OS 
precursors to obtain SiC, more work on the polymerization procedure and 
an understanding of all the process variables is necessary to establish 
standard procedures. 

Although y-ray irradiation can crosslink OS, there needs to be a 
source with a higher dose rate than ^^Co to be practical for PSS. A 
source such as ^^''cs may hasten the crossl inking of OS. 



99 



To produce a large dense monolithic SiC body directly from OS, 
foaming and gas generation should be controlled during crossl inking and 
pyrolysis. In order to control these phenomena, a chemical additive as 
an anti foaming agent needs to be found and/or a new polymer with no 
foaming properties needs to be synthesized. 



CHAPTER III 
SILICON CARBIDE/SILICA COMPOSITES FROM CARBOSILANES AND ALKOXYSILANES 

Introduction 



The significance and advantages of the sol-gel route to produce 
glasses and ceramic materials have been given in Chapter I of this dis- 
sertation. Although the production of silica glass monoliths via the 
sol-gel method is highly significant by itself, the problem of intrinsic 
brittleness of glass may be modified by forming a composite. 

If the strength of a glass were determined solely by its lattice 
cohesive energy, one would predict the glass to be very strong. In all 
cases, glasses are weaker than otherwise expected because flaws concen- 
trate applied stresses. The problem is flaw sensitivity, causing cata- 
strophic failure under applied stresses, and susceptibility to thermal 
shock. Because of these problems, glasses have acquired a reputation 
for mechanical unreliability. 

Many ways to strengthen glasses have been studied, ^8 including the 
following: 

• annealing 

• compressive stresses on the surface via tempering or ion 
exchange 

• dispersion hardening 
fiber reinforcement 

• reduction of flaws 



100 



101 



The strengthening nechanisms and effects at room tenperature for 
glasses may not be the same as the effects at high temperature. In 
metals, conventional precipitation-hardened metal alloys are not ther- 
mally stable and the precipitation may coarsen at elevated temperatures 
well below the softening temperature of the alloy, resulting in a reduc- 
tion of mechanical strength. ^9 In contrast, many ceramics and glasses 
maintain strength and exhibit stability at high temperatures witn low 
density and chemical inertness. 

In an effort to increase strength and toughness, barriers to crack 
propagation in the form of discrete particles or fibers have shown some 
success, ^^ similar to the fine dispersion of second-phase particles long 
used for metallic systems. ^° The strengthening effect in the case of 
metals is attributed to various dislocation impedement mechanisms. ''O 
However, since dislocations do not exist in glasses, particle dispersion 
strengthening of glasses must rely on a different mechanism. Hasselman 
and Fulrath^i observed dispersion hardening effects in certain systems 
of ceramics by dispersion limiting the size of Griffith flaws, thereby 
raising the stress required to initiate or propagate cracks. Lange'72 
has proposed a mechanism that strengthening may occur as a result of a 
line-tension effect due to particles initially pinning a propagating 
crack front and causing a detour similar to that observed for disloca- 
tions. 

Another aspect of dispersion hardening may be the strengthening 
effect of the dispersed particles with a higher elastic modulus than the 
matrix. A fine particle dispersion may also inhibit grain growth or 



102 



crystallization of glass at high temperatures and hence lead to an 
apparent increase in strength. Therefore a distribution of dispersed 
particles on a molecular scale should yield a large inprovement in the 
mechanical properties of a glass. 

Using the techniques described in Chapter II, organosil anes may be 
introduced directly into the monolithic body of a sol-gel derived silica 
glass matrix to obtain a strengthened silica body after appropriate heat 
treatments. The organosilane precursors to SiC are thereby dispersed 
homogeneously in the matrix of silica gel. Upon pyrolysis, the SiOo 
matrix is reinforced by molecularly dispersed SiC particles, thereby 
yielding a molecular composite. Developing procedures and understanding 
the process are the objectives of this work. 

In this chapter, results of the expected hardening of SiOo glass by 
Sic molecular dispersion are presented, as well as the fabrication tech- 
niques used to produce such a body. 

Experimental 
Infiltration of silica gel matrix with carbosilanes 

Silica gel matrices to be impregnated were prepared by hydrolyziny 
tetramethoxysilane (TMOS) at ~90°C with water in TMOS/H2O molar ratio of 
1 mole of TMOS per 4-17 moles of water and adding 0.5-2 w/o of acids or 
1:10 molar ratio of the acid to TMOS. This liquid mixture called a sol, 
was cast in plastic molds and aged at 60-80°C for 24 hours after the 
molds were tightly sealed to prevent liquid evaporation. At the end of 
24 hours some gel monoliths were allowed to age for an additional 24 
hours in the original molds and some gels were transferred to a water 



1U3 



bath for further aging. At the end of approximately three days from the 
beginning, the seals were broken for slow drying in an oven with a 
steady increase in temperature from 80°C to 150°C in 12 hours. The 
typical size of a dried gel was ~20cc. 

The dried gels were decarburized in monolithic forms by heating 
them in air to 500-600°C with a heating rate of --65°C/hour and held at 
500°C for two hours. At this temperature, partial dehydroxyl ation and 
nearly complete decarburation is expected to occur. Tne decarburized 
and partially dehydroxyl i zed gels were stored in a desiccator, after 
cooling to room temperature with ~80°C/hour rate, until the time for 
silane impregnation. 

The organosi 1 anes or carbosilanes used to infiltrate the silica 
gels were polysi 1 astyrene (J-PSS2), vinylic silane oligomer (ViSO), and 
vinylic silane polymer (ViSP). The structures and physical states of 
these silanes have been described in Chapter II. The solid silane 
J-PSS2 was dissolved in toluene or THF to make up ~20 w/o solution. The 
high viscosity liquid ViSP was diluted with THF in 1 g silane/2 ml THF 
for impregnation. 

The impregnation of silanes into the silica gels was carried out by 
placing a monolithic piece of the gel in a silane solution containing 1 
vol% of 3-aminopropyl triethoxysilane as a wetting agent and 5 wt% of 
DCP and soaking it for 2 to 12 hours. For deeper penetration of silanes 
into the gel body, vacuum impregnation for 30 min-6 hours was used. At 
the end of infiltration, the infiltration chamber was brought to atmos- 
pheric pressure and left for 30 min. before transferring it into a 
pyrolysis tube. The processing map is shown in Fig. III-l. 



104 



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105 



The impregnated gels were pyrolyzed in N2 °'" ^" ^'' ^^ various temp- 
eratures with the heating rate typically 50°C/nr. The pyrolyzed SiC/ 
Si02 composite monoliths were stored in a desiccator, after they were 
cooled to room temperature from the pyrolysis temperature, until micro- 
hardness, porosity, surface area, densities, etc. were measured. Micro- 
hardness values were measured by using a Kentron Tester or a Leco Model 
DM-100 with 0.5-2 kg load and a 136°C diamond pyramid indenter. Both 
diagonals of the indentation were measured and the average values were 
used to compute the Diamond Pyramid Number (DPN). Porosity and surface 
areas were determined by using Quantachrome' s Autosorb-6. Densities 
were measured by mercury, water or propylene glycol volume displacement. 
In situ molecular composites of Bulk SiC/$i02 Monoliths . 

For a silica gel matrix, TMOS was hydrolyzed using acidified water 
in 1:4 molar ratio of TMOS/H^O and organic solvents. The water used 
here was previously degassed by inert gas purging to improve the action 
of CFRI IJCP. The acids used in ~2 vol% were lactic acid and formic 
acid. These are strong organic acids and are expected to catalyze the 
hydrolysis of TMOS readily and possibly improve the compatibility of the 
TMOS with organosilanes. Solvents used were isopropyl alcohol, THF, 
toluene, n-butanol , and amyl alcohol in 100-150 v/o of TMOS. Approxi- 
mately 1 v/o of 3-amino propyl triethoxysil ane was added as a wetting 
agent to promote the retention of polysilanes in the matrix. 

After the hydrolysis was complete (~20 min. stirring by a magnetic 
stirrer at room temperature), 0.4-4 w/o of an organosilane (OS) dissolv- 
ed in THF (-1/5 volume of TMOS) with 4-5 wt% DCP (wt% with respect to 



106 



OS) was mixed with the hydrolyzed sol for 20 min. at 30-40°C in a closed 
container. 

The 0S/Si02 sol mixture was cast in Teflon molds. They were sealed 
tightly to prevent any evaporation of solvents and water, followed by 
gellation and aging at ~60°C for 12 hours. Curing was followed with the 
temperature gradually raised to ~150°C in 6 hours. The gels containing 
OS were cured at ~150°C for 6 hours. The temperature was brought down 
to 60°C at the end of 6 hours of curing and the seals were broken for 
slow drying in an inert atmosphere. Drying was continued for two days 
at 60°C in N2 or Ar with a gradual increase of solvent evaporation by 
opening the lids. The typical size of a dried gel was ~20cc. 

Pyrolysis and densi fication were carried out by placing the dried 
composite green bodies in a tube furnace with N2 gas flowing at ~5 ml/ 
min rate and a heating rate of ~20"C/hour up to 150"C and held at 150°C 
for three hours. The temperature was raised to the range of 150-850°C 
with a rate of ~100°C/hour and held at 3bO°C for two hours. The cooling 
rate was also controlled; 200Vhour to 660°C, held at 650°C for one 
hour, 200°C/hour to 300°C, held at 300°C for one hour, then lUU°C/hour 
to room temperature. 

Microhardness, pore volume, surface area, and density were measured 
the same way as for the infiltrated SiC/SiOp composites described in the 
previous section. The processing map for this process is also given in 
Fig. III-l. 
SiC/Si02 molecular composite powder 

A powder form of a SiC/Si02 composite was prepared by mixing 0.6- 
6.3 g PSS, 8-10 ml TMOS in 10-15 ml benzene or toluene, and 0.05-0.8 g 



107 



DCP, in an apparatus shown in Fig. II-3 with N2 gas flushing over the 
solution for two hours before heating began. After two hours of degas- 
ing with N2, the solution was heated to a gentle reflux for 12 hours 
with continuous N2 flowing. At the end of 12 hours, the solution color 
changed from dull gray to brownish yellow. 

The reaction mixture was precipitated in 80/20 dy volume HeOH/HoQ 
solvent and washed with the solvent three times before drying in a 
vacuum oven at 70°C for five hours. After reprecipitation in MeOH, an 
80-90% yield was achieved. 
Characterization of products . 

Fourier-transform IR spectra were taken by using a Nicolet MX-1 
FT-IR Spectrophotometer, SEM micrographs and EDS spectra by a JEOL Model 
JSM-35C, NMR by a Varian XL-100 Nuclear Magnetic Resonance Spectrometer, 
and by Nicolet's High Resolution FT-NMR, BET surface ared and pore size 
distributions by Quantachrome' s Autosorb-6, UV-Vis transmittance by 
Perkin-Elmer's 552 UV-Vis Spectrophotometer, IR transmittance by Perkin- 
Elmer's 283B IR Spectrophotometer, reflectance of visible light by a 
custom built optical microreflectometer. X-ray photoelectron spectra 
were taken by using a Kratos model 800. Hot-stage x-ray diffraction 
patterns in a helium atmosphere were obtained by using a Philips X-ray 
Powder Di ff Tactometer as the sample powder was heated on a platinum 
substrate with HTKIO High Temperature Hot Stage by AP Parr Co. of Graz, 
Austria. Microhardness and fracture toughness of the monolithic compos- 
ites were determined by using a Leco Microhardness Tester and according 
to Antis et al.'^^ Flexural strengths were determined by using an 
Instron Testing machine. 



108 



Variations in experimental conditions, reactant amounts, and chemi- 
cal additives are summarized in Table III-l for the typical composites. 

Results 

A photograph of partially dried OA gel monoliths to be impregnated 
with silanes is shown in Fig. III-2. The pore size distribution of the 
OA gel is given in Fig. III-3. The BET surface areas, mean pore sizes, 
and total pore volumes of typical composites are summarized in Table 
III-2. The volume loading of SiC in the pyrolyzed composites are esti- 
mated from the TGA char yields given in Chapter II for each US and the 
volume OS added to the sol. Approximately 3 vol% SiC was maintained. 

Repeated attempts to produce edge-notched SiOo matrix monolithic 
specimens for fracture toughness were unsuccessful. Attempts to heat 
the composites higher than ~850°C in nitrogen also failed due to the 
foaming of the matrix. 

A photograph of an Si02 monolith strengthened by in situ bulk de- 
composition of an organosilane along with a pure gel matrix and gel with 
OS before thermal treatment is shown in Fig. III-4. 

Infrared transmittance curves of SiC/Si02 glasses with a fused 
silica glass as a control are shown in Fig. Ill-b. A reflectivity curve 
of a bulk SiC molecular composite glass is compared with a blank gel 
glass in Fig. III-6. Ultraviolet-Vis transmittance curves of the compos- 
ites are compared with a blank FA gel glass in Fig. III-7. Figure III-8 
shows IR reflectance curves of FA gel and ViSP/FA gel. An enlargement 
of the band at ~800 cm-1 is shown in Fig. III-9. 

Fourier-transform IR spectra of PSS-10/104 composite powder before 
and after pyrolyzing at 900°C and at 1270°C for three hours are shown in 



109 



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Fig.III-2. Silica Gel Matrices After Aging and Drying 
Prior to Decarburation and Impregnation. 



112 



0.03 




1000 



PORE RADIUS, A 



Fig. III-3. Pore Size Distribution OA Gel Showing no Change in 

Average Pore Sizes After IBO'C and 500° C Heat Treatments 



113 



<4- 
O 


E 


(/) 


3 


3 


> 


'o 

> 


o 


<u 


a. 


o 
a. 


OJ 




■tJ 


1 — 


f= 


■!-> 

o 


o 
<*- 


»> 


T3 




<U 


'^* 


■(-> 


•r- 
T3 


(O 


n3 


3 




O 


o 


O 


Q- 


>1 


c: 


4J 




1/1 

OJ 


•% 


Q 


a; 


0) 

3 
{_ 
I— 


<D 


»a 


o 
><- 




3 




oo 


to 


1— 


4^ 


LU 




CQ 


1/1 




o 




Q. 


CM 


E 
O 


1 


o 


t— « 






C\J 

o 



13 




O 




r- 


>1 


-!-> 


■t-> 


m 


■ f— 


i_ 


CO 


o 


c: 


<D 


OJ 


f— 
3* 


Cl 



•=^ O lD IX) 
r^ ro "^ r~- 









r-^ =S- 3 

oi -^ CM 

I I 



30 O 
CM OO 



o 
a. 



CM CM lo 
r-< ro CM 








CM ys ^ 




CM 


1— 1 LD '^ 


ZO CM 


n 


O -O JT) 


'-n ro 



QJ << 



o 


«« 




ID. 


1/1 


CM r~- 




3 


• • • • 


c 


• r— 


00 'S- CO fVJ 


"3 


"O 


.— 1 1— 1 r-l r-l 


a; 


fO 




2: 


q: 





CO 



1^ o .x> 
t— I .r> <£) 






Lf) 1— 1 in 


00 


CO 


00 LD CO ^o 


00 


■^ 


.— 1 ^ ro >— 1 


. — 1 


<s- 



-O CO ^ .— I 1^ CM LO 

■^ 1^ • i-o ^ •— I 00 

ro CM 00 CM .—I r-~ .— I 



c 

OJ CJ> 
E o 
■(-> 

113 • 
OJ h- 



<u 



0000 

O O iXi o 
CO i»0 r^ 00 



o o 



CO 



f—^ 








000 





Ln 








'-n 


un 


CO 


00 


00 


00 o^ 00 


00 CO 



CJ 



CO 

o 

Q. 

E 
O 

o 

csi 
c 

C_) 



c 










"T" 


3: 











lE 


:r 3r 

















c 


c 


000 


Q- 


O- 


00 


3 


3 <u 3 


u 


1 — OO 00 oo 


OO 


00 


00 


ca 


CO Q. 20 


a. 


OJ .,_ .^ .,- 


• 1— 


■ r- 


2. 


^-^ 


— -* ^"^ ^"^ 


^^.^ 


C7)> S> > 
1 1 1 


> 
1 


> 
1 


1 



00 


n. a. Q. 

00 00 C/1 


00 
to 


<: <: «a: < 


<: 


<: 


=c 


■ r- 


■r— -r- 't— 


a. 


0000 





Ll_ 


U- 


>■ 


>■ => => 


'^ 



114 



J-pss/SiOjgel Composite 



White 120°C 

Blacl( 800°C 

Clear Matrix Pure Gel 






t 



i|i|Mi|l|l|MM|i|i|!mr|Mi|MMMi|MMMMMqMi|MifMiJMi|'ti|'|il'IMi|i|'lM'|i|'l'|ili|ilMl|i|'|i 



'lt!|ll ^l l|llll|if ll| llll^l ll ( ll (l l l ill )ll l ^llll|lfl l |ll ll | l l iy ( l |^ ^ ^ ^ 

' — OK GBS OB OU 001 081 081 Oil 08t 091 On QSl 081 OIL 001 



|i|i|i|i|i|i|r|i| 



Oqc 068 088 QZ2 08Z 
llllllllllllllllllllll 



llllilllllllllllllllllllllllllllllllllllllllllllllllllilllllllllllllllllllllllllllllllllll^^^ 



Fig.III-4. SiC/Si02 Molecular Composite from PSS/SiOo Gel 
Before and After Pyrolysis With a Pure Gel 
Matrix for a Comparison. 



115 



o 






O 

o 
in 

00 




o 


' 




.-1 




in 




o 


< 




r-( 




M 




u 




ra 


- 









u 


^^ 


to 


o 


r-H 


•r^ 


11 


D 


un 




-^ 


^i: 




00 














'Jl 


< 


b 


^ 


o 




[^ 


•'-' 


^-1 


rH 


-o 


a. 


3- 




0^ 


ji 


C/l 













S 




-/ 






o 
o 



o 
. o 






; \ 
i 



I ■ 



1/- 



o 
o 
o 



(U 

o 



T3 

c 
to 



(U 



o 



00 



o 



- c 


— 


TD 


c 


1 


<U 




E 


to 




o 








l^ > 




1^ 


o 

ft3 C 






s- ro 






•4-> ( 


o 




u ••— 


c 




oi on 






ittance Sp 
ed With a 


o 




E -t-" 
00 n3 


o 




C C 


o 




R Tra 
mpreg 



1 % 



O) 



115 







CD 

■a 
<u 
+-> 

(O 

c 

C7) 

O) 

S- 
Q. 

E 



a. 

n3 



0) 



n3 



u- 


o 


o 


o 




CO 


VI 




<u 


■M 


•r- 


<T3 


■l-l 




•r- 


-o 


> 


O) 


•r— 


M 


+-> 


>, 


U 


r— 


tu 


o 


r— 


u 


M- 


>> 


d; 


D- 


IC 





O vC ^ 



117 



10 

ax 



— T ; — 


\ 


















1 "^ 

J 




\ 


\ 








0) 
5C 






1 






\ 






LO 


< 








\ 


\ 






cc 


U. 




' 










U 


<v 








\ 








\^ 


o 


r-^ 


ao 












' "^ 


^\>.^^ 


o 


aj 


01 






[ 








"■ ^V ^N^,^^ 


o 


o 


u 












"^ - ^;r^''"^ 


:». 00 




Q. 






\ 








x:-^^ 




< 






- 


1 








'\ 


^ 


0^ 


a.c 












"\ 




'^ 


c/: o 














^\ 


< 


■H 


•rt u-l 






/ 








\ 


a. 


> 


> 00 






V 








•V 


1 












\ 








-^--' — ^ 


1 












\ 


•< 








1 












/ 


^ 






-_ 


1 












; 








' j 


1 












\ 








\l 












\ 








;) 










■ 


\ 








/ f 












\ 








\ 1 
















_> 




'^ 








— ^ 






\ 


* ^^^^ 


-- 


< 


--■~-^ 
























... ."r-t-ras 












■ 




1 


"*C 


^^ 


\ 1 \ 






















^ 1 ' L__ 



c 



c 
c 

OC 



C: 

O 



X 



1 % 



CO 

o 
o. 

E- 
o 

CJ 

to 



(U 



-a 

c 
re 



CJ 



CO 

■•-> 
<J 
<u 
a. 

CD 



I 
I 



iZ 



118 




o 




CM 






UI 




0) 


C 


1/1 


o 


>1 


L/^ 


^— 




o 




S- 




>i 




c 










: ' 


i- 


CO 


(D 




-!-> 




tl- 




< 


O 




c 


r— 


t-H 


OJ 


— 


o 



CO 

T3 
C 
n3 



C3 



E < 



1T3 



o 

(U 
D. 
00 

(U 
O 

c 

+-> 
(J 



OJ 



I 



CO 
I 



aDUBaoaxjBH % 



119 



ViSP/TMOS 




TMOS gel 



800 , 
v; cm ' 



500 



Fig.III-9. Enlarged Spectra of Fig.III-8 Showing Greater 

Peak Area at --800 cm'' For the SiC/SiO, Composit 



2 Composite 



120 



Figs. III-IO and III-ll. An expanded IR spectrum of J-PSSl/1120 with as 
received J-PSSl for comparison is shown in Fig. III-12. 

Proton NMR spectra of A-PSS-0/927, PSS-10/104, and J-PSSl/1120 are 
given in Figs. III-13, III-14, and III-15. The TGA char yields of the 
powder composites are shown in Fig. III-16. X-ray photoelectron spectra 
of PSS-10/924, A-PSS-0/927, and ViSP/BuOH composites for silicon after 
they have been pyrolyzed are shown in Figs. III-17, III-18, and III-19. 

Scanning electron micrographs of fractured surfaces of ViSO/BuOH, 
ViSP/BuOH, and J-PSS/BuOH are given in Figs. III-20 and III-21. 

The results of a diamond point micronardness test for SiOo gel 
impregnated with silanes are given in Table III-3 and Figs. III-22 and 
III-23 for DPN vs. pyrolysis temperature. An optical micrograph showing 
a diamond pyramid indentation on an impregnated composite is given in 
Fig. III-24. Results for the same test for bulk in situ SiC/SiOo mole- 
cular composites are given in Table III-4 and in Figs. III-25, III-26, 
and III-27. 

The results of fracture toughness measurements directly from an in- 
dentation crack are given in Table III-5. Mean value of 5-10 measure- 
ments were used in the calculation of Kjp. The formula used to convert 
the crack lengths to Kjq described by Antis et al.'^^ -j g given below in 
Eq. III-l. 



Kj(, = 0.016 (^)^^^ ^-372' ^^^"-^ 

where: E = Elastic modulus in pascals 



(c/2)3/2 



P 

H = Microhardness in pascals = 5- 

2(a/2)'' 

P = Indentation load in newtons 



121 




CO 

O 

s- 

D- 

s_ 



C 
fO 

OJ 

s- 
o 

M- 

CO 

o 



r 



4-> 
O 

(U 

c 
tn 

cc 

t 



o 
I 



265 £9> S2£ 

30NV103U3d % 



90-2 S 



122 




t. 
u 



CVJ 



•I— 

CO 

o 

s_ 

>» 

i. 



o 



1^- ° 



1/1 

Ol. 



E 

3 

+-> 

O) 
Q. 



■3ali»a % 



123 



m 



l8-b2 




■-ZZii- 



o 
o 



o 
m 



- E 


t/0 


u 


a. 


•^Jt^ 


1 




•-D 


(/) 




cr 


"O 


oy 


fO 


in CO 


o 


CVJ2 


(NJ 


— ID 
2 


^ 


UJ 


c 


> 


oo 


< 


£X 


O^ 


1 


o 


<+- 


lO 


o 




ra 




&- 




4-> 




U 




cu 




Q. 




I/O 


o 




m 


)— 4 


h- 


1 



0082 6112 SCt'l 

33NV103nd3d % 



ZSZ 



CJ 



124 




QJ 
CD 



o 

s. 



I 

o 
cn 



3 
O 



o 
I 

CO 

oo 
t 



o 

E 

+-> 

u 



o 
+-> 
o 

i- 



I 



125 




CD 



1/5 

</) 

o 

s- 



1 

o 
I 

o 

c 

en 

c 

s 
o 



o 

CM 



CO 

D- 

1 



+-> 
U 
(U 

c 



c 
o 
+-> 
o 

i_ 
c 



O) 



125 




a. 
a. 



CD 



CO 

(/) 
o 

s- 

III 

•1 — 
Ul 

I 
o 

I 
_CVJ 

C3 
I 

CD 



O 



o 



to 
to 



o 

E 

3 
S_ 
-1-1 

<u 

C/) 



c 
o 
■•-» 
o 



127 



<T 


rsi 


<• 


CN 


ON 


o 


a> 


C 


r-H 


o 


1 


O 


1 




1 


m 


0^ 


Ul 


V. 


1 


C/1 




< 





1 


' 1 

1 1 






1 1 


1 I 


; ;' / 


1 


■ 1 1 




1 1 J 




/'' / 




r ' / 




./ y 








-''"^^^ .^^^ 




^^''"^^ ^^^^^'^^ 










/ 




t 


^ — 


f ■ • 




' -^J^ 


""^ 






■'^f 




\J} 




f 









s- 

Ol 

3 

o 

C- 



o 

c 

o 

S- 



(J 

o o 

o -"^ 
O • cvj 

X -r— 
•f— 



o 



o 
o 



-a 









CJ3 



I 



o 
o 



o 



o 

00 



o 



o 






o 



(%)3ripTS3y' • 3n 



128 




SINOOD 



129 




E 
o 

<: 



en 
c 
•I — 
S 
o 



CM 
CTi 
--- <U 

C t: 

I -r- 

Ul ^ 

U~l i. 

Cl. rt3 

I C_^ 

-C 
QJ C 

(/) QJ 
O X3 
CL'i- 

o e> 
</) 

CSJCC 

•r- t/1 

CO c 

■^ o 

C_) (J 



e o 



s_ 


c 


■M 


o 


(-> 




QJ 


+-> 


Q. 


S- 


OO 


o 




Q. 


OO 


O 


D- 


s_ 


X 


Q- 


00 




1 





sxNnoo 



130 



600 



400 



200- 




120 




Fig.III-19. XPS Spectra of SiC/SiO^ Composite ViSP/BuOH 
Proportions of Silicon and Oxygen As Oxide 
and Carbide (top). Oxide and Hydroxide on 
the Surface, Respectively 



131 




Fig.III-20. SEM Fratographs of Bulk in situ Molecular Composite, 
Top: ViS0/Si02, Bottom: J-PSS/Si02 



132 



MJJJf.'^f.Jif" 






v'^^^f 



'■.■ •- ?,v i> 






^v 25KU "X20^0.0* i' :t2B2S%^^^ 









'•f^?:*^'i8!^^-«^i5: 



- ■ . . - . ,;*■ \* ■■-■■.;-. .. 




■ , A ;-' ■-■ . H-; 


;=^V?*f'^■■>^>> I ••-■'■. r^^'^:V.'^ - 


. c - ■ ' -• 




.. ;;.'vi;;. ■; : 


»■ " ■■ -. •■•■;- - ■'. ■!■'.■?■-.'■ '"".;■<■.■■>;•■'.'. ■ 


' "' .- ' 


".■■'-.■■■'*;-■ 


^^isKy' X200B0 


050i*pj.:-l/0U UF^^^ 



Fig.III-21 



Fractured Surfaces of ViSP/Si02 Gel, a Bulk 
SiC/Si02 Molecular Composite Cefore(top) and 
After Pyrolysis at 830°C 



133 



Table III-3. Microhardness of Organosilane Infiltrated 
SiC/Si02 Composites as a Function of Pyrolysis Temperatures 



Composites 
and Control 


DPT^ 


b20°C 


600°C 


700°C 


800°C 


900°C 


OA-ViSO 




139±1 


233±15 


239±14 


669±173 


814±200 


OA-gel 




146±58 


157±1 


174±25 


210±51 


3U5±103 


FA-ViSP 




192±55 


342±75 


35U87 


403±46 


550±180 


FA-gel 




129±64 




171±24 


210±51 


291±67 



134 



1200 



1000 



800 



n OA Gel Control 
OGei/viSO 



600 



400 



20C 







i 



^$' 



{f 



'/:-^---^ 



,0 



() 



[] 



n- 



500 



600 700 800 

Temp, **C 



900 



Fig.ni-22. DPN As a Fuction of Pyrolysis Temperature for 
OA Gel and OA Gel Impregnated l.'ith YiSO 



135 



7oa 



60C- 



50 



400 



300 



20 



00 



/ 

^ 



AFA Gel control 



OGel/viSP 



()/ 



/ 



/ 



() 



/ 



/ 



/ 



<^ / 



o 



A 



/ 



/ 



y 



J^ 



y 



y 



500 600 700 800 900 

Temp^ ° C 

Fig.III-23. DPN As a Fuction of Pyrolysis Temperature for 
FA Gel and ViSP Impregnated FA Gel Composite 



136 




Fig.III-24. Indentation Cracks Formed by Loading of a 

Diamond Indenter at 120C g on FA Cel/ViSP Im- 
pregnated and Pyrolyzed at 8Q0"C, Mag.:200X 
Photo Taken After 3 weeks of Indentation 



137 



Table III-4. Microhardness and Density of ViSP/BuOH, a Bulk 
SiC/Si02 Composite as a Function of Pyrolysis Temperature. 



Temperature (°C) p, g/cc DPN 



600 


0.8±0.1 


250±50 


700 


0.95±0.1 


350.8±134 


800 


1.20+0.1 


500.0±100 


350 


1.45±0.1 


551±169 


900 


1.75±0.1 


366.3±250 



138 




0.5 



600 700 800 
Temp/c 



900 



Fig.III-25. Density as a Function of Pyrolysis Temperature ■for 
Bulk ViSP SiC/Si02 Composite as Compared With Pure 
Silica Gel Matrix 



60C- 



50C- 



139 



/ 



/ 



/ 



/ 



Q_ 
Q 



3oa. 



/ 



/ 



/ 



/() 



/ 



/ 



/ 



<3 



U' 



20 



i 



00 700 800 

Temp , °c 



900 



Fig.III-26. DPN as a Function of Pyrolysis Temoerature 
for ViSP/BuOH Composite 



140 



600 



500 



0.400 
Q 



300 - 



200 







OviSP/siOj 



K3^ 



^>- 




80 



1.0 



-^ 



1.6 



12 1.4 

Density, g/m I 

Fig.III-27. Relationship Between DPN and Density of ViSP 
Bulk SiC/SiO^ Composite and SiO„ Gel Matrix 



.8 



141 









(O 




«s 


■D 




e 






<-} 


^ 




o 




c 


i/i 


k_ 


VI 


TS 




5 






o 


o 




•(— 


OJ 


Q 




O 


O 

.-o 


Ll_ 


t_ 




Ll_ 


T3 
OJ 


«> 


c 


CO 


•^~ 


CO 






o 


O 


1/1 

-l-J 


S_ 


''~ 


o 


en 

O 


X 


E 




O 


• 


O 


LD 


00 


1 


o 






»— 1 


OO 


OJ 


o 


X! 


oo 


<a 





Q. 
O 









o o 



00 
J1 



3 O O O --I 



00 CM 



o ^ 



ITS 


OJ 


po OJ ro 


"=f 


>* 


oo 


Q. 


OJ • 


• • • LO ^ 


• 


• 


• «s- 


s: 


. o 


o o o • • 


o 


o 


o • 




O -H 


+1 +1 4H O O 


+1 


+1 


+1 o 


^ 


+( o 


r^ CTi OJ -H HH 


l^~ 


OJ 


LD +1 


o 


^ a^ 


<o ^ '-n ^ r-< 


CTi 


^ 


"=3- n 



r-l O 



O O O r^ 



fO 


(T) OJ 






OJ ^ 




a. 


+1 +1 


f-H 


1 — 1 


OJ +4 +1 


O^J 


c: 


CO ^ 


+1 


+1 


-H 3D LD 


+1 




• • 


00 


T-H 


'^ • • 


ro 


•* 


OJ o 


• 


• 


. o o 


• 




OJ ^ 


r^ 


LD 


t^ OJ r-H 

OJ o^ r-- 


o^ 


E 


+1 +H 


r-H 


O 


r-H -H -H 




3. 


^ OJ 


1— 1 


(-H 


4H ro oo 


00 




• • 


+ 1 


-H 


T— t • • 


+1 


M 


Tl" O^ 


o 


x> 


O -^ O 


o 


lO 


a^ '.O 


33 


Ol 


T— 1 r~. iD 


1 — 



• 

O 
+1 



1^ 
+1 



ro -^ 
+1 +1 



LO 
+1 

+1 • 

r^ 00 



LO ^ 










+t +1 










( — I . — 1 


-D r^ ^ m ^ 


•n 


«* 


•^ LO 


• • 


+1 -H +1 -H -H 


+1 


+1 


+1 +1 


^H .—1 


-H »3- 3^ (Tl X) 


1-D 


ro 


r^ ro 


"^ ro 


-O ^ ro OJ ^J 


^ 


^ 


ro ro 






T3 
O 



a. 

CD 



OJ 



>> 

o 

O 



3 
3 



<*- 

o 

O) 

3 

I_ 
0} 

■t-> 

-a 
c 
ro 



(U 
o 

c 

<- 



•-n 
o 



lO LO iT) LO 

O O O r-( o 



lO 
«3 



.— I ro OJ o 
OJ OJ 'S- CO 



ro 



o 

LO 



o 

o 
o 



CO ^ OJ >— 1 


'T 


"=3- 


o OJ 


^H OJ P~ Oi 


O^ 


i — 


00 CO 





J3 ro 


O C_3 




0) o 


O 




w) -t- 


O O 




c — 


O O 




OJ -r- 


m CO 




T3 CO 




0) 




r~- r^ 


^ 


>,-c 


dl C_5 t_5 OJ 


Q. 


.— 0) 


010 O) 


E 


1 — on 


o o 


ro 


3 3 


<: o OJ cc 


to 


Ll_ >*- 


o LO CO u- 



t_3 


CJ 




OJ 


• 


o 


3 


o 


•)-> O 


-c O 


4-> 


• r— 


r— OJ 


OJ OJ 


•^■ 


to 


•— 00 


.^ CO 


cn 


^^ 


4- 


ro 




o 


C " 


o - 


c 


•^- 


• r- r— 


CO , — 


• r- 


to 


OJ 


OJ 




CJ 


O CD 


Q. CT1 


o 


^ o 


CO 


to 


U-) 


,— o 


■.- cC 


••- «a: 


• r— 


3 OJ 


5» U_ 


> o 


> 


j3 CC 



^J 


cn 


Q- 


ro 


• 


■o 


o 


0) 




CO 


+ 


3 


a. 0) 


«« 


c^ 3 


>- 


• f^— 


Z 


^H ro 




1 > 


« 


.— 1 


CT 


-^ c 


C 


ro 


-^ 


o<D 


c 


ud E 


c 




o 


II CO 


C_5 


0) 




LiJ -(-> 


* 


o 


to 


31 C 


-^ 


C (U 


i- 


•.- T3 


O 


CO 


2 


3 l- 




OJ 


to 


>>-M 


to 


-Q 4-> 


ro 


0) 


r— 


T3 r- 


ts 


OJ 




-l-J ro 


C7> 


ro 


c 


E <V 


• r— 


•1- > 


C 


-U O 


I_ 


on .Q 


o 


UJ ro 


o 



142 



a = Diamond point indentation diagonal length 
c = Extended crack length 




The results of flexural strengths of these composites obtained by 
using an Instron machine on 3-point bending are given in Table III-6. 
The detailed testing method will be given in the next chapter. 

Hot-stage x-ray powder diffraction patterns for A-PSS-0/927 powder 
on a Pt substrate are shown in Fig. III-28. 



Discussion 
It has been demonstrated that incorporating an SiC phase into a 
pure Si02 glass matrix increases the strength of the composite material. 
However, the art of making large monolithic composites of SiC/SiOo via 
the sol -gel technique is in the successful fabrication of the Si02 glass 
matrix. In this work, it was found that the addition of organosil anes in 
a TMOS sol makes it more difficult to obtain a large monolithic glass. 
Despite the strengthening effect by the SiC phase, the increasing diffi- 
culty of monol ithicity with increasing loading of SiC in the composite 
appears to be a combination of three factors: 

1) the evolved gaseous products from pyrolysis of the organo- 
silane creates flaws and pores (Fig. 11-30), 

2) the SiC phase strongly hinders viscous flow of the glass 
matrix, thus the necking is curtailed (Fig. III-25), and 



143 



Table III-6. Flexural Strengths of 
Organic Derived SiC/SiOo Composites. 



Sampl e 



P, g/cc 



P, N 



Qfiex' MP^ ^flex/p 



12 hr soaked 
OA gel in ViSO 
820°C 


1.26 ± 0.1 


31 ± 5 


4.4 ± 1.5 (3) 


3.5 


Dried OA gel 
150°C 


1.17 ± 0.1 


67 ± 6 
71 ± 6 
62 ± 4 


25 ± 3 (3) 
17 ± 3 
15 ± 3 


21 


HCl gel 

2 hr ViSO soaked 

820°C 


1.75 ± 0.1 


533 ± 12 


73+6 (2) 


42 


HCl gel 
820°C 


1.76 ± 0.1 


408 ± 17 


60 ± 8 (2) 


34 


FA gel soaked 
in ViSO 10 hrs 
800°C 


1.34 ± 0.1 


62 ± 9 


9.2 ± 2 (3) 


7 


OA gel soaked 
in ViSP 6 hrs 


1.59 ± 0.1 


53 ± 8 


25 ± 6 (3) 


16 


OA gel soaked 
in ViSO/ ViSP 
7 hrs 


1.62 ± 0.1 


67 ± 9 


6 ± 2 (2) 


6 


ViSP bulk 

Molecul ar composite 


0.91 + 0.1 


53 ± 10 


5.8 ± 2 (3) 


6.4 



* The numbers in the parentheses denote the number of specimens tested. 



144 





n 900° 


c 






npt 


/ 


SiOz 




Pt 


^ 


I ..j^J\J 


\ 








WV^vr^VVw/V "T;Vr ^-;/V^^^^^.v^^ 


20 


26 
Si02 


^~*vv/*. 


32 

1 


38 
Pt 


44 50 




1 







80 




Fig.III-28. Hot-Stage XRD Powder Patterns of A-PSS-0/927 Composite 



145 



3) a mismatch of CTE between Si02 (0.5 x iO'*^ in/in °C) and 
SiC (4.7 X IJ"^ in/in °C)'^5 increases residual stress (Fig. 
III-24). 
Although these may all be related, the first factor has been discussed 
in Chapter 2 and it seems an inevitable phenomenon for this type of 
organosilane. The second factor was evidenced by densi fication behavior 
as a function of temperature (Fig. III-25). Gel derived silica without 
a SiC phase densified at a faster rate and at lower temperatures. The 
main reason for the sluggish densi fication behavior must be in the 
action of the SiC phase hindering the viscous flow of the glass matrix. 
However, the effect of atmosphere in densi fication behavior may have 
some effect on silica viscous flow, primarily due to water content in 
the atmosphere. 

The third factor that was shown to be harmful to obtaining mono- 
liths could give a positive effect on the strengthening and thermal 
properties of composites by introducing a compressive stress on the 
surface. ''5 

An additional factor which has been repeatedly observed and men- 
tioned previously is chemical reactions between the SiC precursor OS and 
the Si02 precursor gel during the pyrolysis process in a reducing atmos- 
phere. This is an important phenomenon for fabrication of successful 
composites. Gels heavily impregnated with OS always came out like 
crackers: "cracker effect," after the OS/gel composite has been pyro- 
lyzed in nitrogen atmosphere. Although the bodies, many times, general- 
ly remained monolithic, a severe spalling and/or pores were generated. 



146 



As a result, the flexural strengths and densities were significantly de- 
creased. 

The origin of the "cracker effect" may be thought to be from two 
sources. One is the anticipated outgasing by the decomposition of OS 
and elimination of any impurities during the pyrolysis. Tne other 
source is the actual chemical reactions between the reinforcing phase or 
the precursor of the reinforcing phase and the matrix phase as shown by 
Eq. III-2 for an example in the presence of impurities such as residual 
carbon, silicon, oxygen, etc. 

2Si0o(S) + SiC(S) > 3SiO(g) + CO(g) AG = -70Kcal /mol ^7 (III-2) 

^ A 

In this reaction, it can be seen that the glass network is destroyed as 
a result of the reaction. This reaction is farther discussed in the 
next chapter with respect to oxidation of silicon carbide. 

However, as shown in Figs. III-22, III-23, and III-27 and Tables 
III-3, III-4, and III-6, lightly (~2 hours soaking) impregnated compos- 
ites exhibited a significant increase in microhardness , c7^]gx» ^"'^ *^IC 
with little increase in density. 

The determination of Kjq by direct measurements of tne crack length 
has not been universally established. The method is valid only for the 
test specimens which behave normally, i.e. give a well defined radial/ 
median crack system. ^3 Some brittle materials, e.g. glass, have shown 
to give lower Ktq values than values obtained by more standard methods, 
such as single-edged notch beam test.^s These materials have shown slow 
crack growth well after the diamond point indenter is removed from the 
surface of the specimen. Moreover, the method^s has not yet been proven 



147 



as a valid method for composite materials. This is reflected in the 
relatively large uncertainty in Kj^ values (Table III-4). 

The KjQ to p ratio of bulk ViSO reinforced Si02 gel (-1.3) is near- 
ly as large as or larger than hot pressed 20 v/o SiC/borosil icate glass 
composites by Gac et al ."^^ (Kj^/p = 1.6), by Samanta and Musikant'^^ 
(Kjq/p = 1.2), and much larger than sintered chopped Si02 fiber/Si02 
glass composites of Meyer et al .^o (Kjq/p = 0.48). The ratio of ~1.4 
for ViSO/ SiOo 9el composites in this work is also superior to hot 
pressed 42 v/o SiC/Al203 composites by Cutler et al.8i (Kj^/p = 0.80) 
and comparable with zirconia toughened alumina composite by Lange et 
al.82 (Kj^/p ~ 1.4). 

The flexural strength of lightly ViSO impregnated HCl catalyzed 
Si02 gel is increased (Table III-6) to 73 MPa from 60 MPa for the blank 
matrix. This is more than a 20% increase in strength. However, because 
of the statistical nature of strength, the significance of ~20X increase 
is in question and it needs further work. Optimization of the soaking 
time in relation to pore size and distribution is suggested for the max- 
imum increase in af]px i" future studies. 

In order to improve the monol ithicity of the composites, procedures 
in gel matrix preparation may need to be changed. Zarzycki et al.83 
suggested that larger pore sizes promote monol ithicity of gel glasses. 
Yu et al .8'+ supported the suggestion by experimental results. They 
showed that acid catalyzed gels always produce smaller pore sizes than 
base catalyzed gels. They successfully densified base catalyzed gels, 
but not the acid catalyzed ones. This suggests an advantage of the col- 
loidal route to obtain a monolithic glass. 



148 



The acid catalyzed gels used in this work produced cracks and/or 
foaming at temperatures above ~850°C. As shown in Fig. III-3, the ;nean 
pore radius is in the range of 15 A. This is very snail for gases and 
hydroxy! groups to leave the gel structure before pore closure at the 
surface. In addition, larger pore sizes should increase SiC loading by 
OS infiltration. However, it is shown that larger loading of the SiC 
phase is not only unnecessary but also detrimental to the mechanical 
properties of the composite. 

If a successful consolidation of the Si O2 glass matrix monolithic 
composite can be obtained all the way up to full density, the hardness 
and strength of the composite body is expected to be much higher. 

A J-PSS/Si02 bulk molecular composite, as shown in Fig. III-4 after 
pyrolyzing at 800°C for two hours, still maintained a low density of 
1.32±0.1 g/cc. Heating above 800°C usually caused foaming. At this 
stage, to prevent foaming, use of a vacuum or a reactive gas atmosphere 
such as H2 may be helpful. Composites made this way contain no more 
than 4 wt% SiC phase in the gel matrix. 

An SiOo gel film heavily impregnated with a SiC precursor shows 
strong IR absorption in the range from 2.5 ym to 50 ym wavelength region 
(Fig. III-5). In the region of 3.8-4.6 urn, a bulk composite with a low 
level of SiC (~2 w/o) is more IR transmitting than a fully densified 
fused silica control. Above 5 um, they both are IR absorbing. 

The reflectivities of SiC/SiOg bulk composites from J-PSS2/TM0S in 
the visible range (Fig. III-6) are approximately the same. The curves 
above 820 nm are extrapolated by a computer and they may not necessarily 



149 



be accurate. The clear Si02 matrix shows -2% higher reflectivity than 
the black SiC/Si02 composite. The low reflectivity of the gel matrix is 
caused by high transmi ttance and that of the black composite must be 
caused by strong absorption by the black color. 

In the UV-Vis light range, the transmission characteristics of 
these chemically derived materials all absorb UV strongly and transmit 
visible light. The curves for the impregnated composite and the bulk 
SiC dispersed composite are essentially identical (Fig. III-7). The 
spikes at 720, 400, and 340 nm for the composites must be due to the SiC 
phase. 

The infrared reflectance spectrum of a molecular composite ViSP/FA 
gel, as shown in Fig. III-8, is almost identical with pure FA gel, 
except for small absorptions at -2500 cm"^ and -3700 cm"^ for ViSP/FA 
gel caused by adsorbed CO or 002^^ and by adsorbed water, respectively 
(see also Fig. 11-42). The absorption peak at -800 cm"^ is mainly 
attributed to an Si02 tetrahedral response. However, the area under the 
peak for an SiC/Si02 composite is -20% greater than a pure Si02 matrix, 
as shown in Fig. III-9. The unique absorption band for SiC is also in 
the region of -800 cm"^. Hence, the greater peak area for the composite 
must be due to the SiC phase in the composite. 

Figures III-IO and III-ll show a somewhat larger proportion of the 
Si02 phase than the SiC phase in PSS-10/104. This agrees with the in- 
tended proportion, assuming a -30% SiC yield from PSS-10. The propor- 
tion of SiC and Si02 phases are in the range of 40% SiC and 60% Si02 in 
the composite fired at 1270°C. 



150 



Figure III-12 shows the structural changes from J-PSSl to J-PSSl/ 
1120 by the reaction PSS, TMOS, and DCP. It is similar to the cross- 
linked PSS in Figs. 11-20 and 11-21. Proton NMR reveals that A-PSS-0/ 
927 (Fig. III-13) is crosslinked via =Si-0-CH2 by the action of Si(0R)4. 
However, the NMR spectrum of J-PSSl/1120 does not show a C-O-Sin linkage 
(Fig. III-14). Instead, it appears that crossl inking has occurred via 
Si-C-C-Si, as shown for PSS in Fig. 11-47. An NMR spectrum of PSS-10/ 
104 shows, in Fig. III-15, a low level of H^C-O-Si and -CH^-OH. This 
means that Si(0R)4 not only provides the Si02 phase for the composite, 
but also provides crossl inkages for the silane oligomers. Tne TGA char 
yields of these composites are given in Fig. III-16; all show higher 
yield than silane oligomers without Si(0R)4 and DCP. 

An XPS spectrum of PSS-10/924 for Si in Fig. III-17 shows that 
there is -20% SiC with respect to Si species in the SiC/Si02 powder com- 
posite, as intended. The binding energies of peaks all shifted slightly 
towards the high B. E. side due to a differential charging effect. The 
B. E. at 102.75 eV should be for SiC and 108.25 eV should be for Si02 
and/or Si(0H)4 . In A-PSS-0/927 (Fig. III-13), SiC is shown to be about 
one half the Si02 phase. In Fig. III-19, the ViSP/BuOH bulk SiC compos- 
ite shows a small amount of SiC and a large amount of surface water or 
hydroxyl groups^e of the composite, as represented by a peak at ~536 eV. 

Fractured surfaces of SiC/Si02 bulk composites (Figs. III-20 and 
III-21) show rougher and coarser textures than a fracture surface of 
Si02 gel matrix. There are no differences in surface texture between 
80°C dried and 850°C pyrolyzed surfaces of ViSP/BuOH. This is another 



151 



indication that the composite is still highly porous. The SEM micro- 
graph suggests that there is a fracture pattern, as shown by wavy lines, 
in the composites seemingly indicating the direction of the fracture. 

Microhardness values of SiC impregnated gel composites are much 
higher than the gel matrix controls, as shown in Table II 1-3 and Figs. 
III-22 and III-23. Unexpectedly, ViSO exhibits a greater hardening 
effect than ViSP which has the greater char yield. This may be caused 
by two factors: 1) ViSO infiltrated more because of the smaller size of 
the molecules, and 2) ViSO infiltrated more easily and stayed inside the 
pores because of its higher polarity and, hence, had a higher affinity 
to the polar gel matrix by maintaining greater wettability of the gel 
surface than the more nonpolar ViSP. As the hardness increased, the 
diamond indentation crack became more troublesome in measuring the 
indentation sizes, as shown by greater scattering in the DPN values. An 
indentation crack for a ViS?/FA gel composite is shown in Fig. III-24. 

The strengthening effects of the bulk SiC composite ViSP/BuOH (as 
indicated by microhardness) are not as large as the infiltrated compos- 
ites (Fig. III-25 and Table III-4). This must be the result of sluggish 
densification of the Si02 gel matrix in the presence of dispersed SiC. 
Moreover, the hardness value drops sharply with large scattering after 
heat treatment above ~800°C, although with little effect on the density 
(Figs. III-25, III-26, and III-27). This is attributed to the localized 
foaming of the matrix which begins at ~850°C for this type of gel. This 
foaming is thought to occur by entrapped gases in the pores. 

In hot-stage XRD, an SiC/Si02 composite powder showed Si02 crystal- 
line phases at 940°C, but no SiC phase at temperatures up to 14Q0°C. 



152 



This Si02 phase must come from the oxidation of SiC phase, not from the 
crystallization of the SiQ2 glass matrix. 3^ More data on the oxidation 
of SiC will be given in the next chapter. The peaks for cristobalite 
matured slowly with time at 900°C and did not increase further up to 
1400°C in a He atmosphere. This suggests that the SiC phase in the com- 
posite hinders the crystallization of SiOo glass; vice versa may well be 
true also since no SiC crystalline phase is shown at 1400°C in ~1 hour. 
Tnese data are shown in Fig. III-28, along with peaks for the platinum 
substrate. A very intense peak at 29 = 39.4° (2.28 A) may be that of 
tridymite, but it completely disappeared at 1400°C. 

Concl usions 



Sol-gel derived monolithic silica glasses can be reinforced by 
impregnating with SiC by way of an organosilane precursor. The 
reinforcing effect, measured by mi rcohardness, is nearly three times 
greater for the SiC infiltrated composite glass after a heat treatment 
to 900°C than for the matrix under the same condition. Approximately 
100% increase in fracture toughness and a -20% increase in flexural 
strength is achieved with the silane impregnation. 

Incorporating a SiC phase on a bulk scale can be achieved by mole- 
cularly dispersing the SiC from an organosilane precursor in sol-gel 
derived silica. For monol ithicity on a large scale, only 2-4 w/o SiC 
phase by way of an organosilane is allowed. However, it would be poss- 
ible to increase the SiC loading by using a high vapor pressure solvent 
in a high temperature mold with an effective sealing capability. This 



153 



in situ bulk molecular composite gives a reinforcing effect as measured 
by an increase in microhardness values. The presumed transformation of 
the microstructure and physical properties of the SiC/Si02 composite as 
a function of temperature is shown schematically in Fig. III-29. 

An obvious problem in all of these composites is in the densifica- 
tion procedure. Establishing improvements in the sol -gel processing of 
the silica glass matrix is necessary before high performance monolithic 
composites can be made routinely. 

The experimentally observed "cracker effect" needs to be investi- 
gated further to improve composite properties and the fabrication proce- 
dure. 

A SiC/Si02 molecular composite powder with varying amounts of SiC 
phase can be made by mixing an SiC precursor, e.g. an organosilane olig- 
omer, and an Si02 precursor. The Si02 precursor SilOR)^ not only pro- 
vides the Si02 phase in the composite, but also a crosslinking action as 
well. Crystallization of the SiC phase as well as the SiU2 phase is 
suppressed in the composites. 

Although the objectives of developing a procedure and an under- 
standing of the process to produce SiC/Si02 molecular composites have 
been achieved, more work to improve properties is needed for actual 
applications. 



154 




T. "C 



NecKirii 




Liquid 



Fig. III-29. Conceptual Microstructure Transformation and Evolution 
of Physical Properties of Polysilane Dispersed Sol-Gel 
Silica Matrix 



CHAPTER IV 

SILICON CARBIDE/SILICA COMPOSITES FROM 

COMMERCIAL SILICON CARBIDE AND SILICON TETRALKOXIDE 



Introduction 



A keen interest in high performance materials in recent years has 
led to more attention to silicon carbide materials (SiC), as presented 
in the previous chapters. However, the excellent mechanical properties 
and chemical inertness of SiC have not been fully utilized because of 
difficulties in forming and sintering large complex shapes. 

Incorporating SiC in a matrix which has desired properties and can 
easily be formed into desired shapes and sizes may be an answer to the 
problems of forming and sintering SiC. Fiber reinforced composites, 
particularly those incorporating SiC as reinforcement are being increas- 
ingly utilized for their excellent specific properties, i.e. for high 
strength to weight and rigidity to weight ratios. 

The reinforcement of brittle materials with high strength fibers 
can yield composites of very high toughness. This was first demonstrat- 
ed using carbon fibers in glass and glass-ceramics. 88"90 More recently 
the availability of continuous SiC fiber has led to tougher glass and 
glass-ceramic composites^^'^"* which are more resistant to high tempera- 
ture oxidation than the carbon fiber composites. 

Despite its remarkable mechanical properties and chemical stability 
in ambient conditions, gradual oxidation at high temperature limits the 



155 



156 



wider applications of SiC as a truly high performance/hiyti temperature 
material . 

Tnere are several mechanisms known for the oxidation of SiC in the 
temperature range of 1000°-150U°C.95-ioi /\t low oxygen partial pressure 
(Pq ), < ~3 X 10"^ atm at 1400°C, an "active" oxidation occurs due to 
the formation of gaseous products as shown by the Eqs. IV-1 and IV-2 
below. s** 

SiC(S) + 02(g) t SiO(g) + CO(g) (IV-1) 

SiC(S) + 3/202(g) t SiO(g) + C02(g) (IV-2) 

At high Pg , "passive" oxidation due to a protective film of Si02 is 
operative according to the reaction Eqs. IV-3 and IV-4.100 

2 SiC(S) + 302(g) * 2Si02(S) + 2C0(y) (IV-3) 
SiC(S) + 202(g) ^ Si02(S) + C02(g) (IV-4) 

At any case, a gaseous product or products are formed. 

If the latter mechanism of passive Si02 film formation is applic- 
able for SiC oxidation, an intentional coating of the surface of SiC 
material with Si02 glass should help to prevent or passivata the further 
oxidation of SiC. Moreover, the Si02 glass matrix may be used to im- 
prove fabricabil ity of SiC by viscous deformation at much lower tempera- 
tures. Thus the incorporation of Si02 glass into a SiC skeletal struc- 
ture or vice versa by forming a composite has three potential advan- 
tages, 1) improving the fabricabil ity of SiC, 2) reinforcing the matrix, 
and 3) minimizing the oxidation of SiC. 

In spite of these obvious potential advantages, there have been 
only a few studies to investigate how a Si02 matrix can affect the oxi- 
dation kinetics of a SiC/Si02 composite at high temperatures. 



157 



The forming method used by the previous investigators ''5-78 -jri 
making SiC/SiOo composites was the conventional hot pressing technique 
which has great limitations on sizes, shapes and complexities of the 
composite bodies. It is hypothesized that these limitations nay be 
overcome by way of sol-gel processing. As discussed in the previous 
chapters, one of advantages of the sol-gel method is in the simplicity 
and ease of forming a green body. By casting the sol into inexpensive 
or disposable plastic molds followed by aging in an oven, a monolithic 
green body can be produced in almost any shape and size. The previous 
chapters describe methods for preparing a very low volune fraction of 
SiC in a Si02 matrix to form composites. In this chapter, however, a 
process is decribed to achieve a much greater SiC loading in the Si02 
matrix derived from sol-gel technique and commercially available SiC 
fibers and whiskers. 

A lengthy review on ceramic-matrix composites was reported by 
Donald and McMillan. 59 Factors in designing and making fiber-ceramic 
composites were given by Bialoskorski and Konsztowicz. ^^2 jhg factors 
include lengths of fibers, their contents, and types of fibers used for 
reinforcement. Wang and Sutula^^s showed that the detrimental effect 
due to the difference in thermal expansion between fiber and metal mat- 
rix in metal -matrix composites is minimized for short fibers. Lannutti 
and Clarki0'+"i06 showed a potential usefulness of sol-gel derived alum- 
ina in SiC/Al203 composites using whiskers, mats, weaves, and short 
fibers of SiC. Rice et al.if^ showed the effect of interfacial bond 
strength between the fiber and the matrix on mechanical properties. 



158 



Some of important factors affecting mechanical properties of fiber 
reinforced ceramic composites are listed below. 

1. Volume Fractions of Components 

2. Porosity of the Composite 

Lower volume fraction of continuous phase often 
increases porosity 

3. Ultimate Strength of Fiber and Matrix 

4. Interfacial Bonding 

High bond strength between fiber and matrix 
Low bond strength between fiber and matrix 

5. Thermal Expansion Coefficient Match 

Comparability of a thermal property at high temperature 

6. Flaws, Defects, Impurities, etc. 

The objectives of the work in this chapter are 1) developing a pro- 
cedure to produce SiC/Si02 composites using the sol-gel process, and 2) 
understanding the variables affecting mechanical properties. 



Experimental 

Composites of 6-SiC and a-SiC in a pure silica gel matrix were pre- 
pared using Nicalon® and Silar" as the reinforcing filler phase. 

Chopped fibers of Nicalon were pretreated to remove a polyvinyl 
acetate coating by successive washing in ethyl acetate, benzene, and 



* 6-SiC by Nippon Carbon Co., distributed through Dow Corning, Midland, 
MI. 



**a-SiC by ARCO, Greer, SC. 



159 



acetone followed by firing at 400°C in air for two hours. The cleaned 
Nicalon was further chopped in a polypropylene container with alumina 
balls on a vibratory mill for one hour. In addition, continous fibers 
and weaves of Nicalon were cut into the sizes of molds and the polymer 
coating was removed by burning it on a propane burner. Silar" was used 
in its as-received form or after heating in air at 300°C for one hour. 
Scanning electron micrographs of Silar™ and the chopped Nicalon® are 
shown in Fig. IV-1. 

Characteristics of the SiC used are listed in Table IV-1. 

Silica sol was prepared by hydrolyzing tetraethyoxysil ane (TEOS) 
with HCl in ethanol as the solvent in the mole ratio of 1:4:0.5 for 
TEOS:water:alcohol . Ten grams of SiC was mixed with 75 ml of the silica 
sol containing 1-5 ml of glycerol and/or 1-5 ml of formamide as drying 
control chemical additives (OCCA's) for 1-30 min followed by ultrasoni- 
cation before casting in polystyrene molds of various shapes and sizes. 
The various configurations of the composites made in this manner are 
shown in Figs. IV-2 and IV-3. 

The SiC/Si02 sol slurries, after being tightly sealed, were aged in 
an oven at 40-80°C for ~10 hours before slowly drying in air for 5 hours 
at 90°C. The dried green composite bodies were impregnated with Si Go 
sol up to four impregnation-drying cycles followed by a vacuum impregna- 
tion. 

Cylindrical composite bodies for oxidation experiments were cut in- 
to thin wafers using a diamond rotating blade. The wafers were dipped 
into the silica sol twice after each drying in oven. The composite 



160 




® 



Fig.IV-1. SEM Micrographs of Chopped Nicalon Fiber (top) and 
Silar^ Whisker (bottom) 



151 



Table IV-1. Properties of SiC Used in SiC/Si02 Composites, 
Data Provided by the Manufacturers 





NICALON® 


SILAR scg" 


Tensile Strength - MPa 


2000-2520 


689U 


Tensile Modulus - GPa 


180-200 


689 


Elongation - % 


1. 5-2.0 


— 


Density - g/cc 


2.55 


3.17 


Filament Diameter - ym 


13-15 


0.6 


Length - ym 


Infinite 


10-80 


Cross Section 


Round 


Hexagonal 


Type of SiC 


Beta 


Alpha 


Maximum Temperature 


1200°C 


1760°C 


Form 


Chopped or 


Whisker 




Continuous 


Single Crystal 1 ine 




Fiber 





Surface Area - m^/g 
Impurities - w/o 



-20 Oxide 
-IS Free 
Graphite 



Trace of Metals and 
< 3% Oxide 



162 



SiC/SiQ 
SLURRY 




SiC CONTINUOUS 
FIBER / S1O2 




SiC WEAVE/ 
Si Op LAYER 




SiC CHOPPED fiber/ 
Si O2 LAYER ' 



Fig.iy-2. Some of Layouts Used in Fabrication of Nicalon/SiO, 
Gel Matrix Composite ^ 



153 



SiC/Sol-Gel SiO 2 Composites 




Nicalon ® 





Silar 



TM 




Ocm1 Z 3-4"5 



Fig.IV-3. Cast Nicalon /Si02 Gel and Silar"" /5i02 Gel 
Composites of Various Shapes and Sizes. 



164 



wafers were exposed to 1100°C in a box furnace with dry static air and 
sequential oxidation of the wafers were carried out. The samples were 
heated to 1100°C beginning from room temperature each time at a rate of 
~200°C/hr. 

The wafers contained ~55% SiC, -26% Si02, ~15% open pores, and --6% 
closed pores by volume, on the average. The open porosity was measured 
by mercury porosimetry and the closed pores estimated by the theoretical 
density and a subtraction method. Excellent dimensional stability of 
the composite sample wafers was maintained. There was no warpage and 
the shrinkage was < 5 v/o. The mean density was 1.67 g/cm^ which is 
~70% of the theoretical density of the composite wafers. The density 
change was + 0.2 g/cc during the oxidation. 

After each oxidation exposure of the sample, FT-IR spectra were 
taken by using a Nicolet MX-1 FT-IR spectrophotometer in diffuse reflec- 
tance mode. Thermogravimetric analysis in Pq =1 atm. was carried out 
in continuous flowing dry O2 at 1000°C using a DuPont TGA model 1090. 

Some of composites were made using ~ one-third the sol used in the 
cast composite above followed by cold pressing in a steel die (Carver 
Laboratory Press) at -10,000 lbs load for 5 min at 25-80°C followed by 
oven drying and a vacuum impregnation of Si02 sol. The steel die was 
premachined to form notched sample bars for fracture toughness as shown 
in Fig. IV-4 as well as unnotched bars for a flexural strength test. 
Three point flexural strengths were determined by using an Instron Test- 
ing machine and the compressive modulus was measured with an MTS mach- 
ine. A three-point instead of four-point test was used because of the 
specimen sizes. Microhardness was measured with a Kentron Tester using 



Ram 




Mold 



Base 
for C, 



flex 



Fig.IV-4. Steel Mold Used to Cold Press 



Base 
for K,c 

Nicalon®/Si02 Gel 



Silar"'"'^/SiOp Gel Composite Specimen for<^fi 



ex 



and 
and Kj(3 



166 



1-2 kg loads. Porosity was measured usiny mercury jjorosimetry and an 
Autosorb-6. The composite bodies were measured by mercury volume dis- 
placement after each heat treatment. Measurements of fracture toughness 
(Kj(-) were obtained by using the conventional 3-point bend tests on 
notched beams same way as for the 3-point flexural strengths. The in- 
dentation method^s described in Chapter III was also used to determine 
KjQ and compared with the notched beam 3-point bending test. 

A test for thermal shock resistance was carried out by heating cold 
pressed composite bar samples to 800°C in N2 ^'^'^ quenching them in a 
silicone oil bath followed by flexural strength measurements. Effect of 
oxidation on strength of the composites was examined by heating the 
specimen in air at 90U°C for four hours followed by flexural strength 
measurements. 

Transmission electron microscopy (TEM) of Nicalon® and Silar" com- 
posites was used to examine the interfacial region between the carbide 
and the oxide. To make the TEM specimen, thin disk composites were 
formed and cut into ~3 mm diameter disks, then reduced down to ~1 mm 
thickness by an ultrasonic saw. The disks were further polished down to 
~150 urn with successive Carborundum polishing steps. The -150 ym thick 
disks were dimpled by a VCR dimpler. At this stage, the middle portion 
of the samples was ~30-50 ym. Further thinning until a perforation in 
the middle was obtained by using an Ar ion beam of 100 uA current and 6 
KV potential. It typically took ~15-20 hours for the ion milling. 
Transmission electron micrographs were taken using a JEOL, JEM-200CX 



* VCR Group, San Francisco, CA. 



167 



Analytical Electron Microscope with a 20 nm beam size and ~10U sec 
counting times. Thermomechanical analysis (TMA) and differential scan- 
ning calorimetry (DSC) analyses were obtained with a DuPont 1090 Thermal 
Analyzer. 

The effect of OS impregnation into the porous NC and SC was examin- 
ed by soaking the dried specimens into benzene or THF solution (~2 g 
OS/IO ml solvent) for 2-5 hours followed by pyrolysis at 900°C for 2-4 
hours in No and Of-^Q^ measurements with an Instron Testing machine and 
KjQ measurements with a Leco Microhardness Tester. 

Results 



The SEM micrographs of Nicalon® and Silar" in Fig. IV-1 show a ~10 
ym diameter for Nicalon®, and ~1 pm for Silar", and a mean length of 30 
ym for both types of fibers. This gives R - 3 and R ~ 30 for Nicalon*^ 
and Silar", respectively. 

Differential scanning calorimetry of the as-received Nicalon® is 
shown in Fig. IV-5. The exothermic peak at ~320°C at the first heating 
in air is a result of the oxidation of the polyvinyl acetate coating on 
the fiber used for sizing. The peak for oxidation of the coating is ab- 
sent in the second heating. The TGA in Fig. IV-6 agrees well with the 
DSC. A significant weight gain by the oxidation of SiC Nicalon® at 
temperatures above ~750°C is shown in Fig. IV-6. Thermomechanical anal- 
yses (TMA) of a Nicalon composite (NC) and a Silar composite (SC) after 
various heat treatments are shown in Figs. IV-7 and IV-3. Negative ex- 
pansions caused by sintering are shown for all cases. 



168 




100 



Temp^<>C 



400 



® 



600 



700 



Fig.IV-5. DSC Thermograms of as-received Nicalon Heating in Air Twice, 

— Showing the Removal 



The First Heating and second — 

of the Sizing Polyvinylacetate at ~320°C 



69 



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172 



The compressive modulus of cast NC after a 1,00U°C heat treatment, 
yielding a density of 1.79 y/cc, is 571 MPa. Flexural streng-,hs (af^g^)' 
densities, and microhardnesses of the composites are given in Table 
IV-1. Table IV-2 lists the BET surface areas, pore sizes, and the total 
pore volume of NC and SC. Densities as a function of infiltration cycle 
and a function of heat treatment temperature are yiven in Table IV-3. 

Density and afigj. as a function of processing temperature for 1UU% 
chopped fiber NC is given in Table IV-4. The maximum a^^g^ of NC with 
different layouts and processing conditions after 1100°C heat treatments 
for two hours are given in Table IV-5. 

Graphical representations of density vs. firing temperature for 
cast NC and SC are shown in Fig. IV-9. Microhardness (DPN) vs. density 
for cold pressed NC and SC are shown in Fig. IV-10. Greater density and 
DPN are shown for NC than SC. Flexural strength vs. DPN in Fig. IV-11 
also show an exponential type of function. '^f^Qx ^^* ^l^^sity in Fig. 
IV-12 shows a straight line function for both NC and SC. The pressed 
Silar composites (SC) have nearly 2-3x the a-f-jg^ of pressed NC at the 
same density. 

A sample specimen for the fracture toughness represented by Kjq is 
shown in Fig. IV-13. The parameters to calculate Kjq and the resulting 
KjQ values for cold pressed NC and SC are given in Table IV-5. Typical 
load vs. crosshead displacement curves are shown in Figs. IV-14, IV-lb, 
and IV-16. 

Calculation of flexural strengths was done according to Eqn. 
IV-5. 108 



173 



Table IV-1. Three-Point Flexural Strengths, Densities, and 
Microhardness and the Ratio of Strength to Density of NC and SC. 



Specimen a^lg^jMPa p, g/cc Hardness, DPN c^flex'^P 



Cast NC 










100% chopped 
fiber, 800°C 


19+2 (2) 


1.64±0.1 


370±25 


11.6 


100% chopped 

fiber, 140U°C 


83±3 (2) 


1.82±0.2 


660±37 


45.6 


100% chopped 
fiber, ViSP 
infilt., 800°C 


37.2±3 (2) 


1.90±0.1 


467±21 


19.5 


100% chopped 
fiber, ViSP 
infilt., 800°C 


33+1 (2) 


1.23±0.1 


520±52 


26.8 


unidi rectional 
fiber, J-PSSl 
infilt., 1000°C 


9.3+4 (2) 


1.66±0.1 


55U48 


5.6 


unidi rectional 
fiber, 140U°C 


53+6 (2) 
12U±5 (2) 


1.82±0.1 
1.98+0.1 


1051±55 


29.1 
60.6 


Pressed NC 










80°C dried 


12±1 (2) 


1.73±0.2 




6.9 


J-PSSl infilt. 
800°C 


31±2 (2) 


1.85±U.l 




16.8 


800°C and 
quenched to RT 


115±5 (2) 
22±1 (2) 


2.07±0.1 
1.72±0.1 


915±58 


55.5 


900°C 

1100°C 

1400°C 


19±5 (2) 
45±3 (2) 
47+4 (2) 


1.92+0.1 
2.18+0.1 


958±63 


23.4 
21.6 


900°C, 4 hrs 
in Ng 


25 (3) 

29 

40 


1.90 
2.01 
1.97 




13 
14 
20 


900°C, 4 hrs 
in air 


23 (3) 

22 

35 


2.00 
1.88 
1.65 




12 
12 
21 



174 

Table IV-1 (continued). 
Specimen a^-]g^,M°a p, g/cc Hardness, DPN °f]ex'° 



Cast SC 












1000°C 


10±2 


(3) 


1.45±0.2 


378±76 


6.9 


Pressed SC 












80°C dried 


18±3 


(2) 


1.41±0.1 




12.8 


700°C 


31±2 


(2) 


1.80±0.1 




17.2 


800°C 


42±3 


(2) 








900°C 


88±6 


(2) 


1.87±0.1 




47 


950°C 


116±£ 


i (2) 


1.88±0.1 




62 


ViSP infilt. 












820°C 


53+3 


(2) 


1.86+0.1 


403 


28.5 


900°C 


89+7 


(3) 


1.91+0.1 


413+42 


46.6 


950°C 


216+11 (2) 


1.83+0.1 


1284±85 


118.0 


ViSP infilt. 












950°C 


78+6 


(2) 


1.94+0.1 




40.2 


1400°C 


112±8 (2) 


2.08+0.1 




53.8 


800°C quenched 


i25±; 


' (2) 


1.82±0.1 


805±25 


68.7 


to RT 












900°C in N2 
4 hrs 


39 




1.92±0.1 




20 


149 




1.72±0.1 




87 




117 




1.76±0.1 




66 




88 




1.71±0.1 




51 




68 




1.87±0.1 




36 


900°C in air 


45 




1.88±0.1 




24 


4 hrs 


101 




1.91+0.1 




53 




95 




1.75±0.1 




54 



The numbers in parentheses denote the number of specimens tested. 



175 



Table IV-2. The BET Surface Areas, Mean Pore Sizes, 
and Total Pore Volumes of NC and SC. 





Heat T 


rec 


itment 


BET 




Me. 


an Pore 


Total Pore 


Composite 


Temp 


• > 


°C 


Area, m^ 


Vl 


S 


ize, A 


Vol . , cc/g 


Cast NC 


300 


in 


ai r 


141 






16.2 


0.114 




800 


in 


N2 


45 






16 


0.036 




1500 


ir 


1 ai r 


0.86 






163 


0.0071 


ViSP infilt. 


800 


in 


Ng 


99 






13 


0.063 


NC 






b. 












Pressed NC 


300 
900 






73 

4.6 






14 
68 


0.0517 
0.0155 


Cast SC 


300 
850 






184 
18 






17 
14 


0.16 
0.13 


Pressed SC 


300 
850 






143 
5.8 






18 
140 


0.12 
0.040 


Pressed/ 


800 






29 






33 


0.049 


ViSP infilt. 


















SC 



















176 



Table IV-3. Density Changes as Function of Processing 
Temperature and Sol Impregnation Cycles for Cast NC and SC, 



Dipping Cycle After 
500°C for 2 Hours 



1 

2 

3 

4 

5 



Density, g/cc ±0.1 



NC 


SC 


1.44 


1.11 


1.56 


1.29 


1.70 


1.41 


1.81 


1.52 


1.86 


1.59 


1.87 


1.62 



Temperature, °C 
After 2 Dipping Cycles 

500 

900 

950 
1100 
1400 



Density, g/cc ±0.1 



NC 


SC 


1.72 


1.41 


1.80 


1.49 


1.88 


1.54 


1.98 


1.62 


2.11 


1.73 



Two specimens were used for each cycle and each temperature. 



177 



Table IV-4. Density and Flexural Strengths of Cast 
NC 100% Chopped Fiber as Function of Processing Temperature. 



Temperature, °C 



p, g/cc ± 0.1 


Gfiex' ^'P^ 


^flex 


1.52 


15±4 


9.26 


1.71 


17±3 


9.94 


1.83 


20±4 


10.9 


1.95 


27±5 


13.8 


2.06 


83+3 


40.3 



80 

500 

900 

1100 

1400 



Two specimens were tested for each tenperature. 



178 



Table IV-5. Maxinum a^ig^ °^ ^^ °^ Different 
Layouts and after Processing at 1100°C or 1400°C, 



Composite "^^flex' ^^^^ 



cast 100% 27±5 (3) 

chopped fiber 83±3 (1400°C) (2) 



bidirectional weave/ 15±3 (4) 

chopped fiber 



continuous fiber/ 25+4 (2) 

chopped fiber 120±5 (1400°C) (2) 



cold pressed 115±12 (3' 

100% chopped fiber 



The numbers in parentheses denote the numbers of specimens tested. 



179 



2.4 ^ 




Fig.IV-9. Density as a Function of Firing Temperature of 
Cast Nicalon®/Si02 Gel Composites and Silar"""/ 
SiO„ Gel Composites 



180 



1200 - 



1000 - 




1.4 



t 



2.0 



1.6 P :^A^^-9 

Fig.IV-lO.Microhardness as aFunction of Density for Cast Nicalon / 
SiOp Gel and Silar'^^/SiOj Gel Composites 



181 



1400 



1200 



^ 



1000- 



800 
Z 

600 



400 



200 




20 



40 



60 



C^lex ,A^P 



L. 
00 



120 14(5 ^^ rio 



Fig.IV-n. Microhardness vs Flexural Strenth of Cold Pressed 
NC and SC 



182 




2.2 



f. 3/c. 



-®/ 



Fig.IV-12. Flexural Strengths of Pressed Nicalon /SiO Gel and 
Silar"'"/Si02 Gel Composites As a Function of Density 



183 




Molded Notch 



4.4Cm 



Fig.IV-13. Schematic of Notch-Beam Test in Three-Point Bending for 
Fracture Toughness of Pressed NC and SC 



to 
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Cast, 1100°C 
Nicalon® Continuous fib 





_L 



J. 



0.05 



0.1 



0.3 



Fig 



0.15 0.2 0.25 
Crosshead Disp. , inch 
IV-14. Load vs Crosshead Displacement of ViSP Impregnated^ SC A/ter^ 
Pyrolysis at 



900 



*C{top) and Micalorf Cont.. Fiber/Si02 Gel Comp .( bottom) 



186 




sqx 'peal 



187 




sq-[ ' pecq 



188 



0,1,, = 3/2 ^^Q^^^-^P^") ^ (IV-b) 

^^ (width) (thickness)^ 

Calculation of Kjq is based on Bansal and Duckworth ^o^ and shown by Equ. 

IV-6. 

1/2 
K,- = ^ ^^^,. ^ Y (IV-6) 

^^ b Vl^ 

where M = applied bending moment at fracture = P x S 

Y = dimensionless paraneter which depends on a/w and type of 
loading 

= Aq + A^ (a/w) + A2(a/w)2 + A3(a/w)3 + A4(a/w)'^ 

for s/w = 8, Aq = 1.96, A^ = -2.73, A2 = 13.66, A3 = -23.98, and A4 = 

25.22. 

A fracture toughness measurement by a Vicker indentation method^3 

has been described in Chapter III. The results of the crack lengths, 

hardness, and Kjq are summarized in Table IV-7. The relation between 

crack length and Kjq is as follows, although the validity of this method 

has not yet been fully proven for composite materials. 



H = ^—y and K.p = 0.016 (|)^/^ ^^ (IV-1) 

2(a/2)^ ^^ " (c/2)-^^'^ 



where: H = hardness in Pa 
P = load in newtons 

a= mean diameter of Vicker indentation 
c= mean diameter of the extended crack due to loading 
E = elastic modulus estimated using the rule of mixture 
The TEM micrographs of cast NC and SC before and after heat treat- 
ment at 900°C are shown in Figs. IV-17, IV-18, IV-19, and IV-20. 



189 



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190 





Fig.IV-17. Bright Field TEM Micrograph of NC After 80°C Drying 
at 59KX Mag. (top) and STEM Micrograph of the Above 
at 300X Mag. (bottom) 



191 




Fig.IV-18. Bright Field TEM Micrographs of Cast 
Heat Treatment in N^. Top: 50KX Mag. 



NC After 900 C 
Bottom: 1 KX Mag, 



Darker Contrast = fiber and Lighter Cont. = Silica 



192 




Fig.IV-19. Bright Field TEM Micrographs of SC After Drying at 
80°C, Top: 50KX Mag. and Bottom: 200KX Mag. 



193 




Fig.IV-20. 



Bright 
in No. 



Field TEM Micrographs of SC After Heating at 900' 
Top: 50KX Mag. and Bottom: 200KX Mag.. 



194 



Figure IV-21 shows SEM micrographs of a surface of cast NC and SC 
after drying at 80°C. Figure IV-ZZ shows a polished surface of NC after 
1100°C for two hours in N2. Figure IV-23 shows fractured surfaces of 
unidirectional continuous Nicalon® fiber in Si02 matrix showing dislodg- 
ing of the interfaces and of random 100% chopped fiber NC. 

The extent of 1100°C dry static air oxidation of SiC in the Si02 
composite was quantitatively estimated using reflectance values (R) of 
FT-IR and the base-line method. -i*' The base-line method was applied to 
compensate for the fluctuation in IR response of the instrument. A few 
typical FT-IR spectra are given in Figs. IV-24, IV-25, IV-25, IV-27, and 
IV-28. The IR bands to estimate R values of each species are tabulated 
in Table IV-8. Reflectances at -789 cm"^ (R^.), -1086 cm'^ (R ), and 
-1157 cm"-'- (R^) are due to vibrational motions of the Si-C bond, 
vitreous Si-0 bond, and cristobalite Si-0, respecti vely . ^ ^ ^"i 13 Plots 
of time vs. R values were made and are shown in Figs. IV-29, IV-3U, 
IV-31, and IV-32. 

In the composites, the total amount of vitreous silica is shown to 
be relatively constant throughout the oxidative heat treatments. This 
means that the oxidation of SiC in Si02 glass produces predominantly a 
crystalline Si02, namely, cristobalite. This was demonstrated by Figs. 
IV-24, IV-25, IV-26, IV-27, and IV-28. 

Thus the normalization of R's by R„ should compensate R due to sam- 
ple shapes, sizes, and surface morphology after each heat treatment. 

The results of isothermal TGA carried out in dry static oxygen are 
given in Figs. IV-33, IV-34, IV-35, and IV-36. The weight gains are 
tabulated and given in Table IV-9. 



195 




Fig.IV-21. SEM Micrographs of As-Cast NC (top) and As-Cast SC 
(bottom) After Drying at 80°C. 



196 




Fig.IV-22. SEM Micrographs of Polished Surface of NC After 1100°C 
Heating in N2. 



197 




Fig.IV-23. Fractographs of NC, Top: Uniaxial fiber after 1C00°C 
Bottom: Random chopped fibers after 1100''C 



198 




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203 



Table IV-8. Summary of FT-IR Spectra for a Sequential 

Oxidation of SiC and SiC/Sol-Gel Derived SiOo Conposite in 

Dry Static Air at 1100°C, ~55 v/o SiC, -25 v/o SiO^. 



Exposure 



~"l " _ 1 _ 

Time, hr "^x, cm"-^ R„ ^'g 

A 


, cm"-'- 


«y 


^c, cm-1 


«c 


R,/Rg 


R,/Ry 


Nicalon® (s-SiC) Composite 








1167 0.004 


1086 


0.10 


789 


U.072 


0.04 


0.72 


5 1257 0.22 


1114 


0.15 


800 


0.14 


1.44 


1.0 


10 1260 0.14 


1118 


0.30 


798 


0.14 


0.47 


0.53 


16 1265 0.49 


1120 


0.31 


794 


0.10 


2.0 


0.78 


Raw Nicalon® (s-SiC) Parti ci 


jlate 












undetectable 






790 


0.1 





oo 


11 1256 0.075 


1104 


0.16 


798 


0.069 


0.47 


0.43 


26 1260 0.15 


1112 


0.17 


816 


0.10 


0.88 


0.59 



Silar" (a-SiC) Composite 






1126 


0.01 


1087 


0.13 


792 


0.15 


0.08 


1.16 


5 


1246 


0.17 


1126 


0.15 


790 


0.20 


1.13 


1.37 


10 


1250 


0.34 


1124 


0.22 


788 


0.38 


1.55 


1.72 


16 


1256 


0.27 


1126 


0.14 


788 


0.21 


1.93 


1.50 


26 


1260 


0.31 


1126 


0.11 


736 


0.17 


2.82 


1.55 



Raw Silar" (a-SiC) Whisker 






1259 


0.029 


1152 


0.038 


805 


0.293 


>0.77 


>7.7 


11 


1282 


0.054 


1156 


0.015 


806 


0.423 


>3.6 


28 


26 


1248 


0.142 


1144 


<0.01 


764 


0.249 


>14 


>25 



204 



2.4 



O Rx/Rg 

ARc/Rg 



NC 




10 15 20 30 

Time, hr 



40 



Fig.r/-29. FT-IR Rflectance of NC As a Function of the Exposure 

Time in Air at nOO°C for a Carbide and a Cristobal ite 



205 




10 15 20 

Time, hr 



Fig.IV-30. FT-IR Reflectance of Raw Nicalon As a Function of Time 
Under Oxidative Exposure at 1100°C in Air 



206 



X X 

Composite <^ R^ A^ 

0.5 L S^9 O 

Raw Siiari'^ r^^ p 



SC 




10 15 20 

Time, hr 



25 



30 



Fig.IV-31. FT-IR Reflectance of Silar'7Si02 Composite and 

Raw Silar^'^As a Function of Time Under Oxidative 
Exposure 



207 



R 




10 15 

Time, hr 



20 



25 



Fig.IV-32. FT-IR Reflectance of Silar /SiOo Composite and Raw Sil 
As a Function of Time Under Oxidative Exposure, After 
Normal ization 



208 



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A heat treatment of the composites up to 1500°C in air showed no 
visible change in shape or physical state. Hov;ever, heating the compos- 
ite bodies in vacuum (~10"^ torn) at 1400°C resulted in a disintegration 
of the composite body and an evolved gas resembling H2S. This must have 
been caused by some type of chemcial reactions between the two phases 
which has been noted in Chapter III. 

Discussion 



Several differences between the fibrous SiC Nicalon® and Silar" can 
be listed: 

1. phase: Nicalon®; 6-SiC 

Silar"; a-SiC 

2. dimension: Nicalon® is ~10 times greater in 

diameter and length 

3. crystalline form: Nicalon® is polycrystall ine fiber 

Silar™ is single crystal whisker 
Extremely large particle sized SiC phases i.e. continuous fibers 
and weaves as shown in Fig. IV-2 showed much lower density and strength 
in cast NC. Fracture energy, however, is expected to be higher as shown 
in Fig. IV-16. 

The polyvinyl acetate sizing on Nicalon® could be removed by heat- 
ing in air at ~35U°C as shown in Figs. IV-5 and IV-6. As shown in Fig. 
I\/-6, heating the fiber in air above ~700°C caused weight gain by SiOn 
formation. 

Shrinkage due to heating begins at ~200°C and accelerated above 
~700°C (Figs. IV-7 and IV-8). At 700°C -10 m linear shrinkage is shown 



214 



for both NC and SC. Greater slope for SC after drying at 80°C is shown. 
Considering the initial sample sizes, NC shows a greater densification 
rate of 0.48% linear shrinkage vs. 0.26% linear shrinkage of SC in ten 
hours. This must be partly due to the greater initial density of NC in- 
dicative of better compaction of Nicalon® and suggests a better compati- 
bility between the silica sol and the Nicalon® fiber phase than that be- 
tween the Silar whiskers and the matrix. This agrees with the results 
in Fig. IV-9. 

The three-point flexural strengths given in Table IV-1 indicate 
that there is general agreement among of-^Qy^, p, and hardness. However 
the large variations in Of-^Q^^ from sample to sample are probably due to 
the number of variables involved in the sample preparation; 1) particle 
size and degree of agglomeration, 2) sol concentration and wettability 
of the sol, and 3) the amount of surface oxide on the starting SiC. 
Infiltrations of SiC precursor polymers in the porous composite bodies 
result in a little increase in Of^Q^ (Table IV-2) and Kjq (Table IV-7). 
The density values suggest that very little silane is infiltrated under 
the conditions used. 

A linear type of relationship between a^-ig^ and p is shown in Fig. 
IV-12. This means that one may need only density measurements to pre- 
dict flexural strengths. Extrapolations based on the Fig. IV-12, 
assuming a theoretical density of the composites of 60 vol% SiC:2.4 g/cc 
for NC and 2.8 g/cc for SC, result in a theoretical a^^g^ °^ ^^^ ^^^^ ^°^ 
NC and 425 MPa for SC. These maximum theoretical values are still lower 
than the experimental Of-^^^ of ~700 MPa obtained by Prewo and Brennan^^ 



215 



by hot pressing borosilicate glass and continuous unidirectional Nicalon 
fibers. This shows the difference between random fiber orientation and 
uniaxial fiber orientation in af]g^. In the uniaxial fiber composites, 
the fiber length to diameter ratio, R = «. The fact that a SC gives 
much greater CTf]gx f^^^t be attributed to the R values, R ~ 3 for NC and 
R ~ 10 for SC. This supports the theory^os that the reinforcement 
effect of a composite is a function of the "aspect ratio," R. This is 
consistent with results of Sambell et al.sa.iu for carbon fiber/glass 
matrices composites. 

No attempts, other than incorporating unidirectional continuous 
fiber, were made to orient the chopped fiber in the Si02 matrix. The 
resulting composites are thus expected to have randomly oriented fibers 
and are treated as being isotropic. The volume fraction of SiC was kept 
at a constant level of ~60 v/o for cast composites and ~70 v/o for 
pressed composites. Variations in SiC loading were limited by the inher- 
ent porous matrix of gel derived silica and monol ithicity of the compos- 
ite body. The former limits the SiC volume loading toward a lower level 
and the latter tends toward higher volume loading. This forces the SiC 
loading to be at approximately a constant level (~60 v/o) of the satura- 
tion to maintain monol ithicity. 

Based on a small number of specimens used for a testing of thermal 
shock resistance of the composites, both NC and SC have good thermal 
shock resistance between 800°C and room temperature. 

Figure IV-9 shows that heating the composites at atmospheric press- 
ure up to 1400°C in a reducing atmosphere did not yield the theoretical 



216 



density. This imposes a major problem in densifying the composites. 
Heating the composites in water vapor may nasten the densif ication pro- 
cess but oxidation of SiC by the water may be a problem. A vacuum heat- 
ing to 1400°C was shown to be detrimental to these composites by disin- 
tegrating the composites. This is thought to be a result of chemical 
reaction between Si02 and SiC according to Lq. IV-?.^"^' '''^ 

2 SiU^lS) + SiC(S) t 3 SiO(g) + CO(g) aG = -70 Kcal/mole (IV-7) 
This is probably the main cause of the "cracker effect" observed for the 
organosi Iane/Si02 gel composites in Chapter III as well. 

Silicon carbide ceramics have modest fracture toughness, expressed 
by the critical stress-intensity factor, Kjq : ~3 MPa •m^/^. ^ ^5. 116 jhg 
NC and SC, however, have Kjq values as high as -8 MPa.m^'^ for SC and ~6 
MPa«m^'^ for NC (Table IV-6). These are remarkably high toughness 
values considering the modest Jfig^ of these composites. This high 
toughness must be attributed to the action of yielding fibers which 
bridged the crack surface. This is evidenced in the load vs. displace- 
ment curves in Figs. IV-14. 

Among the various mechanical properties, fracture toughness as 
expressed by Kjq is considered as important a parameter in design con- 
sideration as the various strengths. This parameter provides a measure 
of tolerance of the material to presence of discontinuities. In some 
applications, the fracture toughness is even more important than 
strength. The fracture toughness relates to fracture strength, impact 
resistance, and thermal shock and fatigue behavior. ii^ This implies 
that NC and SC are potential candidates for heat-engine materials. 



217 



The generally lower mechanical properties of cold pressed NC than 
tnose of cold pressed SC are attributed to the lower lenyth to dianeter 
of fiber ratio (R).ii2 For Micalon^ R ~ 3 and R ~ 10 for Silar". 

Figure IV-14 shows increases of fracture energy represented by the 
area under the curves by silane impregnation into SC and into uniaxially 
laminated NC. The initial sharp increase of load must be attributed to 
the elastic behavior of the composites. Until the crack is propagated 
to the fiber, no load increase is shown. This is the stage of the weak 
matrix cracking. Then the sharp increase in load is the stage of fiber 
controlled crack propagation and finally to failure. 

Composites fired at 1400°C show a typical load vs. crosshead dis- 
placement of a brittle material. The toughness curves shown in Fig. 
IV-16 are essentially the same as those for flexural strength (Fig. 
IV-15). 

One of factor listed as important in mechanical properties of fiber 
reinforced ceramic matrix composites is the interfacial bonding between 
the fiber and the matrix. ^^'^ As a result of chemical reactions between 
the fiber and the matrix, the bond strength between the two phases can 
be increased and this can cause fiber degradation. Rice et al.^O'' have 
shown that this fiber-matrix interaction can be controlled by a third 
chemical reagent as an additive to the matrix or as fiber coatings. 
Characterization of fibers, matrix, and fiber-matrix interactions thus 
requires microanalytical techniques because the scale over the critical 
processes occur is quite small; 2-3 ym across the interface. 

Figure IV-17 shows NC after 80°C drying. The 59,000X TEM micro- 
graph reveals that the particle-like gel matrix has a distinct interface 



218 



with Nicalon*^ fiber. It shows no matrix-fiber interaction at 80°C. The 
STEM micrograph at 300X, however, shows debondings or cracks in the 
interfaces. These cracks must have developed during the thinning pro- 
cesses. This means that the bond between fiber and matrix is relatively 
weak in cast SiC/Si02 composites. The composite after heating to 90U°C 
in nitrogen shows the gel matrix is no longer particle-like as shown in 
Fig. IV-18 at 50KX. The interface is certainly less distinct than the 
composite at 80°C. This may be an indication of tne fiber matrix inter- 
action. At IKX, cracks in the matrix are still present. 

Figure IV-19 shows a SC sample after 80°C drying. The TEM micro- 
graphs at 50KX shows a random orientation of Silar" whiskers bonded by 
SiOp gel matrix with a distinct interfacial boundary. This is also 
shown by the micrograph at 200KX. After 900°C heating, the extension of 
the matrix into the whisker is not clear (Fig. IV-20). The whiskers 
show a distinct columnar subgrain structure representing crytal 1 ographic 
mi sorientations. This kind of planar stacking disorder is common in 
SiC, partly because of the very low stacking fault energy {OUUl} of a 
planes. ^13 Corresponding selected area diffraction patterns exhibit 
spotty rings for the a-SiC single crystal. 

The as cast surface of NC after a 110U°C heat treatment. Fig. 
IV-21, shows the random orientation and distribution of the fibers in 
the Si02 matrix. Figure IV-22 shows that the fibers and Silar" whiskers 
are well bonded by the gel matrix after drying. 

Scanning electron micrographs of fractured surfaces of NC with uni- 
directional and random chopped fibers show fiber-pull-out as indicated 



219 



by the rough surfaces. This either disagrees with the work of Bender et 
al.^^3 or an extensive matrix-fiber interaction may not occur at temper- 
atures below ~1200°C. This may be another advantage of forming a SiC/ 
Si02 composite via sol-gel process. In the work of Bender et al.,ii9 
they used hot-pressing to form Nicalon®/Si02 composite at temperatures 
well above 1300°C. At this temperature, the fiber-matrix interaction is 
expected to be quite extensive resulting in a significant reduction of 
mechanical properties. 1^9 

Rather large scattering of data points are seen in Table IV-1, 
Figs. IV-9, IV-IU, IV-11, and IV-12. This is caused by the difficulty 
of using one identical sample throughout the thermal treatments as well 
as the characterization steps. Even each composite sample prepared 
under identical experimental conditions tended to be different after an 
identical thermal history. This is because there are other variables 
present such as nonuniform property and distribution of the starting 
materials, uneven distribution of flaws, impurities, etc. which affect 
the properties. 

As given in Table IV-1 and shown in Fig. IV-37, the mean Of^Q^ 
after treatment at 900°C in N2 for four hours is 92 ± 26 MPa for SC and 
31 ± 6 MPa for NC. Under same condition but in air, af^g^ ^^ ^^ * 23 
MPa for SC and 27 ± 5 MPa for NC. Both NC and SC exhibited -14% and 
~13% reduction respectively in the strength by the oxidative exposure. 
However, within the uncertainty range, this reduction in strength may 
not be significant. Clark et dl.120 showed a reduction of Nicalon® 
tensile strength from ~1900 MPa to ~8b0 MPa after treatment at 1UUL)°C in 



220 



110 










100 
90 


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NC 


in N2 


NC in Air 


30 




R 








20 . 




WM 


__ 



licalon/SiOp Gel and Silar/Si02 Gel Composites 



Fig.IV-37. Flexural Strengths of SC and NC After Exposed 
to Air or Nitrogen at 900 °C for 5 hours. 



221 



wet air for twelve hours. This ~80% reduction in strength may not De 
used in direct comparison with HC because of the differences in the ex- 
perimental conditions and the nature of the sample. NC was heated in 
dry air at 900°C for four hours while the Nicalon fiber of Clark et 
al.i20 \^as heated at 1000°C for twelve hours. Moreover, the testing 
method is tensile vs. bending. However, it may be inferred that Nicalon 
fibers in NC are more resistant to oxidation than the naked Nicalon 
fiber in a similar environment. 

Figure IV-38 compares flexural strengths of NC and SC with some 
other ceramic composites of comparable composition fabricated by other 
investigators as given by the reference numbers. Comparisons here are 
based on the most recent literature and not necessarily on the highest 
values available in all the literature. The purpose of these compari- 
sons is not to show the superiority of the NC and SC of this work over 
other composites of similar composition but to show tnat NC and SC are 
not inferior or, rather, tnat they are comparable in some cases. These 
other ceramic composites were mostly fabricated by hot pressing techni- 
que. A direct comparison shows that SC and NC are generally inferior to 
other composites but superior to tnose of Gac et al.''^ and Wilson and 
Breit.121 

The lower values for NC and SC are due to lower densities. When 
this is taken into account by calculating the a^ig^/p ratio, the values 
for NC and SC become comparable with other ceramic composites as shown 
in Fig. IV-39. Figure IV-39 compares the ratio of flexural strength to 
density of NC and SC with various other reinforced ceramic composites by 



222 



600 



500 



400 

(T , 

flejj 

MPa 
300 



200 



SiC fiber/ 
Vycor 



THIS WORK 

Cold 
Pressed 



Cast MC 




(92) 



(79) 
30v/o SiC/ 
tnullite 



(78) 
SiC/boro- 
(121) silicate 
SiC whisker/ glass 
Si02 




^^ 



100 



CERAMIC COMPOSITES 



Fig.IV-38. Flexural Strengths of NC and SC as Compared with Some Other 
Ceramic Matrix Composites 



223 



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224 



other investigators as given by the reference numbers. Cast or cold- 
pressed NC and SC of this work are comparable with other composites of 
similar composition fabricated by hot pressing technique. The Nicalon'^/ 
Vycor® composites of Prewo and Brennan92 show exceptionally high a^-, / 
p. This exceptional property is, in part, resulted from the uniaxial 
orientation of continuous Nicalon fiber and two-step process i.e. pre- 
form the composite using wet traditional ceramic processing then hot 
pressing. Lange et al.82 showed 1340 MPa in Of-^^^ value for their ZrO^/ 
AI2O3 in a four step process. On the other hand, NC and SC were fabri- 
cated in a one step process. 

Figure IV-40 compares Kj^ values of iNC and SC (Table IV-6) with 
other comparable composites. Fracture toughness of values of NC and SC 
compare well with those of other composites. Tne Kjq of NC and SC are 
generally superior to the other composites with exception of those of 
Prewo and Brennan9i and Ricei23 for SiC fiber/glass matrix composites. 
The SiC fibers in these exceptional composites were unidi rectional ly 
preformed followed by hot pressing procedure while NC and SC are random- 
ly oriented short fibers and whiskers without hot pressing. Even more 
exciting comparison of NC and SC with other ceramic composites for the 
ratio of Kj^ to density is shown in Fig. IV-41. 

It has now been shown that mechanical properties of NC and SC are 
comparable or superior to many other "the state-of-the-art" ceramic com- 
posites of similar composition made by more conventional, tedious, and 
restricted techniques. Moreover, fabrication of SiC using sol-gel tech- 
nique still has a room for improvements as discussed previously. 

The composites NC and SC upon heating at 1100°C in air caused the 
oxidation of SiC as shown in Figs. IV-24 through IV-28. At 11U0°C, both 



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227 



the vitreous SiOo cind Cristobal ite can be formed as products of SiC 
oxiddtion. Cristobal ite formation rapidly progressed as indicated by a 
reflection band at -1170 cm"^ (R^) for both s-SiC (Nicalon) and a-SiC 
(Silar). Figures IV-29 through IV-32 illustrate this phenomenon more 
vividly. 

A large amount of vitreous SiOo matrix in the composite should sup- 

1100°C 
press the forward reaction of the oxidation, a-SiC — t^ *■ vitreous 

^2 

SiOj. As a result, formation of cristobalite over vitreous Si02 

predominates. This was supported by an XRD analysis of the heat treated 

composites. 

For Nicalon® (B-SiC) in silica matrix, the rate of cristobalite 
formation, which is a direct result of oxidation of SiC at 1100°C, 
rapidly increased up to ~5 hours then decreased up to ~10 hours (Fig. 
IV-29). The rate increased again after ~10 hours and slowed down after 
~15 hours. On the other hand, the raw Nicalon (no matrix) showed a 
continuous increase of the rate of cristobalite formation. 

A similar but less dramatic behavior is shown for the Silar" 
(a-SiC) composite. Silica crystal formation rate for the raw Silar" 
continued to increase, however (Figs. IV-31 and IV-32). 

These complex oxidation rates of SiC may be explained by works of 
Pampuch and coworkers. 12"+. 125 They showed a growth of a ternary inter- 
mediate species Si-O-C which could not be etched by HF. 

In Nical on/Si O2 composite, a rapid increase of R^/'^q ^^^^ ^^"^^ ^P 
to ~5 hours may be an indicative of the initial oxidation of SiC. As 
the Si02 film from the initial oxidation of SiC passivates or decrease 



228 



the oxidation rate, the reaction of SiC + SiUg to yield SiOC may take 
place after ~5 hours of an oxidative heat treatment of Nicalon/SiO^ 
composite. In that case, the amount of Si02 and SiC would decrease as 
long as the reaction rate to give SiOC is appreciable. However after 
~10 hours of oxidative exposure, the rate of Si02 growth increases and 
slows down after ~15 hours. This may be the period of a reaction 

SiOC +02^ Si02 + CO, (IV-8) 

thus increasing Ry/Rr, and decreasing R./R^. The increase in R^/R^ after 

Ay i_ y ■- y 

~15 hours, however, may be due to a development of microcracks in the 
matrix by evolution of gaseous products, e.g. CO, CO2. This is also a 
cause of the continuous oxidation of SiC after ~2U hours, although at a 
slower rate. 

The TGA data of SiC and the composites as given in Table IV-8 
indicate that the weight gain due to SiOo formation per unit weight of 
sample is much smaller for SiC in a SiOo glass matrix than the raw SiC. 
In 5 hours, raw Nicalon® (s-SiC) gained more than twice the % wt than 
did the Nicalon fiber in a glass matrix. The Silar" (a-SiC) whiskers in 
a glass matrix are nearly BOX more resistant to oxidation than raw 
Silar" within 12 hours of oxidative heat treatment. 

The ^J^n vs. time plot for Silar" given in Fig. IV-32 does not 
show the downward curve as does Nicalon®. However the Ry/Rr, values due 

X y 

to oxidation of Silar" somewhat flattened between 10-15 hours of oxida- 
tive exposure. This may mean that the formation of the SiOC intermedi- 
ate phase in Silar™ is not as extensive as in Nicalon. The period of 
SiOC also came somewhat later (~10 hours) in Silar" than in Nicalon® (~5 



229 



hours). It appears that a little longer time is needed for Silar" to be 
passivated than for Nicalon. This may be due to the greater SiOg con- 
tent in as-received Nicalon® than in as-received Silar™. 

Therefore, the ternary intermediate SiOC phase is formed between 
~5-10 hours for Nicalon® and ~10-15 hours for Silar" in porous Si02 
glass matrix. However this SiOC is formation was not detected for raw 
SiC samples (Figs. IV-30 and IV-31). This can be explained by the phys- 
ical properties of the SiC materials. Silar" is a single crystal a-SiC 
with greater density (3.2 g/cm-^) of hexagonal crystal and greater pur- 
ity, while Nicalon® is a microcrystal 1 ine 6-SiC which has a cubic form 
with a density of 2.6 g/cm^^ and substantial amount of graphitic carbon, 
Si02, and Si as impurities. These impurities should act as oxygen scav- 
engers hindering oxygen diffusion to reach SiC unlike al uminosil icates, 
where the aluminosil icate impurity increased the oxidation rate of hot- 
pressed SiC by enhancing oxygen diffusion. 126 

It appears that the ternary intermediate SiOC exists between SiC 
and Si02 interface^os and unless the SiOo blocks oxygen diffusion 
completely, the reaction of SiOC + O2 t Si02 + CO should continue until 
the Si02 filni brings complete passivation. It seems more likely then 
that oxygen diffusion is the rate limiting step rather than CO removal. 

The trend of deceleration of the oxidation rate of SiC in the gel 
derived glass matrix as shown in Figs. IV-29, IV-31, and IV-32 after ~5- 
10 hours has been explained by the formation of a SiOC ternary inter- 
mediate phase. This similar trend was also observed by Suzuki et al.i27 
in their study of a-SiC oxidation. They observed a break in the curve 



230 



of increasing SiO^ layer thickness after ~6 hours of oxidative exposure, 
but no explanation as to what caused the break was given. Leei23 also 
has observed the similar trend of a break in the increasing oxidation of 
SiC with time in oxygen atmosphere, but he treated this anomaly as an 
artifact despite the repeated observations. Again a similar behavior of 
SiC oxidation in a gel derived alumina matrix was observed by LaTorre et 
al.^29 jhey 3]so gave no explanation for this behavior. 

Conclusion 



Low density SiC/SiOo composites can be prepared easily by a sol -gel 
method without hot pressing. Complex shapes can be easily cast using 
SiC particles dispersed in the SiO? sol. 

The ease of forming and fabrication, stability at high temperature 
with a moderate flexural strength, high fracture toughness and high 
hardness have been demonstrated. It was further demonstrated that NC 
and SC have attractive strength-to-weight ratios and even more attrac- 
tive toughness-to-weight ratios, good oxidation resistance, and good 
thermal shock resistance, yielding a potential ability to withstand 
severe environment. Achieving the full density of the composites will 
undoubtedly improve the properties. Considering the low density of the 
composites tested (p ~ 1.9 g/cc), the toughness observed (Kjq ~ 6-8 
MPa«m-'-'^) is truly remarkable. The Kjp of NC and SC is superior to many 
other ceramic composites and the Kj^/p values of NC and SC are superior 
to all but highly oriented hot-pressed samples. 

The high porosity of NC and SC is due to the large inherent shrink- 
age in the silica gel matrix. The flexural strength was shown to be 



231 



directly related to the density of the composite body. The lower den- 
sity and lower strengths of the Silar" system indicate that the Si02 sol 
is less compatible with a-SiC Silar" than with B-SiC Nicalon®. This 
must be because of the amorphous nature of Nicalon and the large amount 
of Si02 and graphite present in Nicalon®. 

Although an improvement in density and strength can be achieved by 
multiple impregnation of Si02 sol into the micropores of the composite 
bodies by dipping, there is a limit on the improvement since the impreg- 
nation is confined to a the surface region after ~4 cycles. 

Impregnation of the porous NC and SC with polysilanes, SiC precur- 
sors, followed by a pyrolysis show a little improvement in the flexural 
strength after each cycle. 

Nicalon® and Silar" SiC have less oxidation in air at 1100°C when 
SiC is incorporated in a sol-gel derived Si02 glass matrix than when raw 
SiC is exposed to sane condition. The oxidation kinetics of SiC in Si02 
matrix (p = 1.8 g/ml ) after heating in air at 1100°C is 3-5 times slower 
than the pure SiC under same conditions. However due to the residual 
porosity, complete immunity of SiC oxidation in the matrix in the begin- 
ning was not observed. 

Both Nicalon® and Silar" in a Si02 porous glass matrix may be oxi- 
dized via formation of a ternary intermediate SiOC phase. The oxidation 
reaction mechanism of SiC in porous Si02 glass matrix appears to be com- 
plex and requires further work to understand completely. Silar" (a-SiC) 
is more prone to oxidation than less pure Nicalon (g-SiC). This may be 
due to the larger surface area of Silar" than Nicalon®. 



232 



In order to improve the oxidation resistance of SiC further, it is 
necessary to reduce the porosity of the composite and prevent natrix 
damage after a prolonged oxidative exposure at high temperatures. 

Some immediate applications of this material may be in fusion power 
reactors. Silicon carbide material because of its low plasma contamina- 
tion, low induced radioactivity, capability of high operation tempera- 
ture, and relatively abundant raw material supply, may be a leadiny can- 
didate in the plasma chamber of a fusion machine. A formed SiC/Si02 
bulk composite may be used directly or as a metallic part coated with 
this composite. As mentioned in this chapter, coating of a material is 
expected to be as simple and easy as casting. 

Another area of application may be in high temperature radomes. 
Requirements for high temperature radomes demand a material with a high 
toughness, good thermal shock resistance, high resistance to rain and 
sand erosion, and light weight. The composites described in this 
chapter should meet these requirements. 

The high values of Kjq/p and/or Of^^^/p of these materials may be 
especially significant in space applications such as materials for ad- 
vanced spacecrafts, space transport systems, and large scale antenna 
arrays, which require high thermal performance and light weight. The NC 
and SC should provide significant improvement in thermal stability and 
mechanical properties and be capable of in-space processing. It should 
be possible to control the elastic modulus and stiffness of such space 
structures and their damping capacity by varying the volume fraction of 
fibers or whiskers in the composites. 



CHAPTER V 
OTHER CHEMICALLY DERIVED CERAMIC COMPOSITES 

Introduction 



The advantages of chemically derived ceramics and composites have 
been presented in the previous chapters. In this chapter, processing of 
several other chemically derived ceramic composites of mullite fibers 
and gel derived silica, alumina powder and polysil astyrene, SiC/SiC com- 
posites of Nicalon® and Silar" with polysi 1 astyrene, and molecular com- 
posite powders of SiC/TiC and SiC/Al203 are presented. 

The mullite fiber is commercially readily available as Nextel™* 
which is derived from a sol-gel process. A monolithic composite incor- 
porating Nextel" into a sol-gel derived silica matrix is attempted simi- 
lar to the processing of the SiC/Si02 composite in decribed Chapter IV. 

Using the processing techniques of SiC from PSS as described in 
Chapter II, it is anticipated that a SiC fiber/SiC matrix composite can 
be made. A successful fabrication of SiC/SiC composite should yield 
many desired properties. 

Wei and Secher^^s ^ere able to increase Of^Q^^ from 500 MPa to 680 
MPa and Kjq from 4 to 6 MPa.m^/^ by incorporating fine TiC particles in- 
to a SiC matrix followed by hot pressing at 2000°C. Yajima et al.i2 
have obtained SiC/TiC powder by mixing their polycarbosil ana with a 
Ti(0R)4 followed by pyrolysis to above 1500°C. Hence, it is hoped to 
synthesize a SiC/TiC molecular composite powder using polysilastyrene to 



4= Manufactured and provided by 3M Company, St. Paul, MN. 

233 



234 



demonstrate that the desired ceramic naterial can be obtained by 
chemical processing. 

Silicon carbide/AlgOj composites have been made by Wei and Becher-'i 
and by Cutler et al.^^ by hot pressing mixed powder of SiC and A1203. In 
this chapter, a novel way of synthesizing SiC/Al203 composite powder at 
the molecular level is demonstrated by mixing the precursors of SiC and 
AI2O3. 

Experimental 

Nextel" 312 continuous fiber which has a mullite composition was 
provided by 3M Co. Nextel" fibers were cut and stacked unidi rectional ly 
with a desired thickness in a rectangular plastic mold. The fiber stack 
was filled with silica sol as described in Chapter 4. After covering 
and sealing the mold, it was aged, dried, and infiltrated with the sol 
4X as for Nicalon® composites in Chapter 4. After an appropriate heat 
treatment in air, density and flexural strengths were measured the same 
way as described for Nicalon® and Silar" composites in Chapter 4. 

Alumina/silicon carbide composites were prepared by mixing 3 g of 
Baikalox CR6* fine powder with 5 ml of 20 w/v PSS solution in benzene 
followed by cold pressing in the steel die as shown in Fig. IV-4. The 
cold pressed green bodies were pyrolyzed to convert PSS to SiC at 80U°C. 

Silicon carbide/silicon carbide composites were prepared by mixing 
3 g of chopped Nicalon® or 3 g of Silar" with 5 ml of 20 w/v PSS solu- 
tion in benzene and cold pressing as above. 



* a-Al203 with 1 ym particle size by Baikowski International. 



^35 



A molecular composite of SiC/TiC powder was prepared by mixing 10 g 
J-PSS2, 0.82 g DCP, and 8.5 ml of titaniun isopropoxide in 60 ml of tol- 
uene under ^2 at 110°C in an apparatus shown in Fig. II-3 and a gentle 
reflux condition for 14 hours. The solution had a dark green color at 
the end of 14 hours. The product was precipitated out by ethanol and 
washed with ethanol three times before vacuum oven drying. The dried 
and agglomerated product had a bright orange color. Nearly 100% yield 
was achieved. 

A molecular composite of SiC/Al203 was prepared in a similar way to 
SiC/TiC above but 0.75 9 DCP and 9 ml of aluminum sec-butoxide with 
refluxing for 22 hours under No. The product was precipitated out by 
ethanol and washed twice with methanol followed by vacuum oven drying at 
150°C. The dried agglomerate appeared dull white in color and ~70% 
yield was achieved. 

Fourier-transform IR spectra before and after pyrolysis of the com- 
posites were taken oy a ilicolet MX-1. Hot stage XRD by a Philips Elec- 
tronic Automated X-ray Powder Di ff Tactometer with HTKIO Hot Stage was 
used to obtain x-ray diffraction patterns of the molecular composites. 
The composite powders were introduced on a platinum substrate using a 
polymeric glue. Resistive heating was applied to the substrate in helium 
atmosphere. 

Results 

Mechanical properties of Nextel" composites and Al203/SiC compos- 
ites are given in Table V-1. Char yields of molecular composites pow- 
ders are given in Table V-2. Energy dispersive spectra of PSS/A1(0R)3 



236 



Table V-1. Mechanical Properties and Densities of Nextel'"/Si02 
Composite, SiC/SiC Composites, and Alumina/SiC Composites. 



Composites-t 


afiex, MPa 


Nextel (60)/ 
SiOg, 900°C 


29±3 


Silar (70)/ 
PSS, 800°C 


15±2 


Nicalon (70)/ 
PSS, 800°C 


30+3 


AloOo (75)/ 
PSS, 900°C 


22±2 



KjQ, MPa m^/^ 



0.43±1 
2.41±1 



DPN 



p, g/cc 



1.92±0.1 



434±28 1.60±0.1 



55U32 1.65±0.1 



1.87±0.1 



* Three specimens used for each composite. The numbers in parentheses 
are voli in the pyrolyzed composite. 



237 



Table V-2. Char Yields of Molecular Composite Powders 



Composite 



Wt% 
SiC 



% Weight Char Yield 
at 940°C in N2 for 14 Hours 



PSS/Ti(0R)4 
PSS/A1(0R)3 



30 
35 



25±6 
55±7 



238 



and PSS/Ti(0R)4 before pyrolysis are shown in Fiys. V-1 and V-2. They 
show the constituent elements Si; Al and Si; Ti respectively. 

Fourier-transform IR spectra of PSS/A1(0R)3 after drying and after 
pyrolysis at 900°C are shown in Figs. V-3 and V-4. Figure V-5 shows 
FT-IR spectra of PSS/Ti(0R)4 composite before and after pyrolyzing at 
940°C. 

Scanning electron micrographs of PSS/A1(0R)3 and PSS/Ti(0R)4 after 
pyrolyzing at 940°C in N2 and hand-grinding are shown in Fig. V-6. They 
all show agglomerates. 

Hot-stage XRD of the pyrolyzed molecular composite powders are 
shown in Figs. V-7 through V-15. 

Di scussion 



Nextel" fiber is not a strong fiber. The tensile strength data, 
provided by the manufacturer, 3M Co. is ~10 MPa. The ofigx °^ Nextel/ 
Si02 composites after multiple sol impregnation and 90U°C sintering in 
Table V-1 is modest after considering the relatively low tensile 
strength of the fiber. The significance here is that Nextel" fibers can 
be formed into monolithic structural composites if necessary. A higher 
temperature heat treatment should improve density and, thus, mechanical 
properties. 

The SiC/SiC composites made from Silar"/PSS and Nicalon®/PSS after 
pyrolyzing at 800°C yield a somewhat low af-jg^ mainly because of the 
high porosity. Sintering of these composites to full density is expect- 
ed to be difficult by the reasons given in Chapter 2. 



239 




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WAVENUMBERS (cm"') 



Fig.V-5. FT-IR Spectra of PSS/Ti (OR) After Drying at 60'C and Pyrolyzing 
at 940'C ^ 



244 




Fig.V-6. SEM Micrographs of PSS/AU0R)3 (top) and PSS/Ti(0R)4 (bottom) 
After Pyrolyzing at 940°C 



245 









Pt 
1 


Pt 




..-.Jj. 


1 


1 


/^V-^jW^X/^Vv-yy-J V^ 


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Fig.V-7. Hot Stage XRD Pattern SiC/Al203 Composite, Top: at RT After 
Pyrolysis at 940''C in N2, Bottom: at ISSO^C after 10 min. 



246 



Pt 



Pt 




20 



26 



l^A\,,A<^AVl«A^yvW' 



32 



38 



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50 



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50 



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32 



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3S 



Pt 



Pt 





50 



50 



56 



68 



29 



74 



80 



Fig.V-8. Hot-Stage XRD Pattern of SiC/Al203 Composite, 
Top: at ISSO^C after 50 min. 
Bottom: at 1400°C after 12 min. 



247 



(0 

C 
0) 



26 



Pt 



48 



56 



Pt 




38 



Ttr 



Pt 



Pt 




64 



72 



IB 



50 





80 



Fig.V-9. Hot-Stage XRD Pattern of SiC/Al203 Composite, Top: at 1400 C 
After 30 min.. Bottom: at RT After 45 min. at 1400 C 



248 




iq V-10. Hot-Stage XRD Pattern of PSS/Ti(0R)4 Composite, Top: at RT 
After Pyrolysis at 940°C, Bottom: at 900°C After 10 min. 



249 



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L 

80 




Fig.V-n. Hot-Stage XRD Pattern of PSS/Ti(OR), Composite, Top: At 900"c 
After 40 min,, Bottom: At IIOO'C After 10 min. 



250 




Fig.V-12. Hot-Stage XRD Pattern of PSS/Ti(0R)4 Composite, Top: at 12G0*C 
After 13 min.. Bottom: at IZOCC after 30 min. 



25' 




20 



26 



32 



Pt 



44 



50 



^0 



56 



62 



68 



74 



80 



(0 

C 
0) 



Pt 



20 



26 



32 



38 



44 



50 



Pt A 




50 



56 



62 



68 



74 



80 



20 



ng.V-13. Hot-Stage XRD Pattern of PSS/Ti(0R)4 Composite, Top: at 1300'C 
After 12 min.. Bottom: at 1300*C After 30 min. 



252 



20 



^ 



(0 

c 



TieOii 



26 



56 



PC 



32 



38 



Pt 




62 



68 



74 



50 



80 




Fig.V-14. Hot-Stage XRD Pattern of PSS/Ti(OR), Composite, Top: at 1360*C 
After IC min.. Bottom: at 135C°C After 30 min. 



253 



Pt 




20 



> 



26 



32 



Pt 



50 




(050 

Z 
Ul 

I- 

z 



56 



62 



68 



74 



80 



TigO^^ or 



SiC or Tie 



Pt 




Pt 



16 



Tt 



Pt 




32 



40 



■^ 



40 



48 



20 



64 



72 



Fig.V-15. Hot-Stage XRD Pattern of PSS/Ti(0R)4 Composite, Top: at 1350 
After 1 hr.. Bottom: at RT After ISSO^C for 1 hr. 



254 



Alumina/SiC composites formed from a-Al203 fine powder and PSS fol- 
lowing pyrolysis at 80U°C show similar mechanical properties and densi- 
ties to SiC/SiC composites above. Since SiC phase is the matrix, the 
sintering may be difficult. More work is needed to show the strength as 
a function of heat treatment temperature. 

The char yields of PSS/Ti(0R)4 and PSS/A1(0R)3 composites are mod- 
est to low. These are crosslinked polymers which are insoluble in com- 
mon solvents. Yajima et al.^^ postulated that their polycarbosil ane is 
crosslinked via Si-O-Ti. In the PSS system, because of the absence of 
hydrogen directly bonded to Si, the crossl inking was expected to occur 
via Si-CH2-0-Ti and via Si-CH2-0-Al for PSS/A1(0R)3. 

These are only speculations and there is no way to know the precise 
nature of the chemical reaction between PSS and A1(0R)3 and between PSS 
and Ti(0R)4 without further studies. Further study varying the reactant 
ratio and amount of the CFRI should be helpful. 

Energy dispersive spectra in Figs. V-l and V-2 show the constituent 
metallic elements in the green polymers. 

Fourier-transform spectra of PSS/A1(0R)3 in Figs. V-3 and V-16 
resemble the spectrum of PSS more than that of A1(0R)3. However, the 
Al-O-Al vibrational mode is shown at -1100 cm"^. The band at 1100 cm" 
became more prominent after the composite was pyrolyzed to 900°C (Fig. 
\/-4). A reflectance response representing the Si-C stretching mode is 
shown at -800 cm"^ as a step. In Fig. V-4, it can be seen that the 
pyrolyzed PSS/A1(0R)3 composite indeed has Al-O-Al and Si-C as expected. 

Figure V-5 shows the FT-IR spectra for PSS/Ti(0R)4. Most of Si-CH3 
features in PSS are lost after crossl inking with TilOR)^. After 



255 



Al (Sec-Butox)3 



o 

z 

o 



UJ 

on 




J- pss/AKSec-Butox), 



5600 



4400 




800 



200 



3200 2000 1400 

WAVENUMBERS, cm"' 

Fig.V-15. FT-IR Spectra of PSS, A1(0R)3, and PSS/A1(0R)3 After Drying 



255 



pyrolyzing at 940°C, the bands for Ti02 (-750 cm"^) and SiC (-800 cm"^) 
are shown. The conversion of Ti02 phase to TiC phase is only possible 
at temperatures above HOO^C. Yajima et al.^^ showed the transformation 
of Ti02 to TiC at 1700°C. Thus at 940°C, the composite must consist of 
SiC and Ti02 and that is what is shown in Fig. V-5. This is documented 
by hot-stage XRD as shown in Fig. V-10. A small amount of titanium 
dioxide, actually TigO^^^^, is formed below at ~950°C. As temperature of 
the composite in He is increased, the XRD intensity for TigOii did not 
increase. The peak at 54.3° 2e (1.69 A) cannot be identified. It is 
not SiC, TiC, Si02, AI2O2, nor Ti02. As temperature increased to 1100°C 
(Fig. V-11), a new and sharp peak appeared at 52.5 29 (1.74 A). This 
peak may be an artifact or an intermediate species because it dis- 
appeared immediately after raising the temperature. At a temperature of 
1200°C, the peaks for TigO^^ are diminished (Fig. V-12). A further 
diminishing is observed at 1300°C and after 30 min at 1300°C, they all 
disappeared. At 1360°C, even platinum peaks are diminished. After one 
hour at 1360°C and cooling to room temperature, the small peak at 36° 29 
representing TigOii came back. However the position of 36° 29 is the 
position for SiC as well as TiC. The true identity of the peak at 36° 
29 is not clear. The reason why the TigOii crystalline phase dis- 
appeared at 1300°C is unknown since the m.p. of Ti02 is ~1840°C. This 
may be the temperature at where the transformation from Ti02 to TiC and 
solid state reactions among the components might begin. 

It appears that the crystallization of TigO^]^ is somewhat 
suppressed as discussed in Chapter 3 for Si02/SiC composites. This 



257 



phenomenon is even more strongly demonstrated in the PSS/A1(0R)2 
composite. Figures V-7 through V-9 show that no crystalline phase is 
present in the SiC/Al203 composite at a temperature as high as 1400°C 
for 40 min except the peaks for the Pt substrate. 

Figure V-6 shows that the SiC/Al203 and SiC/TiOg composites powders 
are agglomerated. 

Conclusion 



Composites of Nextel" fiber/SiOg gel and Al 2O3 powder/SiC from PSS 
can be formed with modest flexural strengths. Silicon carbide/silicon 
carbide composites can also be formed easily and also have modest mech- 
anical properties. Temperatures higher than ~900°C for sintering should 
improve the density as well as the mechanical properties of these com- 
posites. 

The moleciilar composites of SiC/Al203 and SiC/Ti02 are formed from 
chemicals of metal -organic derivatives. They appear to have a submicron 
particle size and suppress crystallization of each phase in the compos- 
ite at temperatures up to -1400°C. A higher temperature than 1400°C is 
required to transform SiC/Ti02 to SiC/TiC. 

The primary and common problem associated with the physical and 
mechanical properties of the composites produced herein is the densifi- 
cation. Use of sintering aids, a controlled atmosphere with higher 
temperature capability should help the problem. 

It has been further demonstrated in this chapter that the potential 
to produce desired ceramic materials, which are difficult to obtain by 



258 



the conventional processing, via chemical process is enormous. However, 
this enormous potential can only be made useful through a continueo 
understanding of the chemical processes involved and elimination of 
uncontrolled porosity. 

The significance of this chapter should be in tne concepts rather 
than a production of exciting mechanical properties of these composites. 
The concept of mixing two precursors of a composite at a molecular level 
can be applied to many other systems. Testings for mechanical proper- 
ties of these composites must be followed after a successful synthesis 
of a molecular composite is made. 



CHAPTER VI 
CONCLUSIONS AND RECOMMENDATIONS 

Silicon carbide material can be made from silicon and carbon con- 
taining polymers in a relatively easy manner. These precursors can be 
used to strengthen the brittle materials such as glass by forming com- 
posites. These composites show increased strength and fracture tough- 
ness even with the low temperature heat treatments and low densities 
(1.7-2.1 g/cc). 

It becomes rather clear that there are variety of functional groups 
which may be introduced to the polysilane backbone. A systematic study 
of how these groups are related to char yield of silicon carbide, vis- 
cosity, green density, chared density, susceptibility to oxidation in 
green state and in the pyrolyzed state and crossl inkabi 1 ity is recom- 
mended. 

Changing the porosity and pore size of silica gel matrices could be 
used to vary the amount and depth of SiC precursor silanes and thereby 
improve the control over the physical and mechanical and/or optical 
properties of silane impregnated silica glass monoliths. To achieve 
this control, the behavior of pure silica gel monoliths with respect to 
pore collapse and densi fication must be better understood. 

Use of a nonpolar hiyh vapor pressure solvent should help to in- 
crease the loading of the SiC phase in monolithic SiC/SiU2 molecular 
bulk components. However, the increasing difficulty of maintaining 



259 



260 



monol ithicity of the composites with increased loading of the SiC pre- 
cursor needs to be investigated as well. 

The chemical reaction between the SiC precursor and the Si02 gel 
matrix is a serious problem and needs to be understood. A systematic 
study of the reaction should help to produce better composites and is 
highly recommended. 

Although the Nicalon/Si02 and Silar/Si02 composites have a remark- 
able strength-to-density ratio and toughness-to-density ratio, further 
study to consolidate the bodies is needed. Complete consolidation may 
be possible using a controlled atmosphere high temperature furnace. 

The molecular composite powders of SiC/Al203, SiC/Si02, and SiC/ 
Ti02 should be hot pressed to determine their mechanical behavior. The 
SiC/Ti02 composite may be converted to SiC/TiC composite by heating at 
temperatures of ~1700°C. A TEM study to examine how the two phases are 
arranged would be helpful in optimizing the process. This idea of mol- 
ecular composite may also be applied to other systems such as Zr02 and 

Tinted and tempered glasses may be made by an impregnation of the 
SiC precursors within the ultraporous Si02 gel glass matrix. This glass 
should produce very high strength and toughness with good IR absorption 
characteristics which may conserve heat for example in cooking wares. 
The high temperature limit of these wares may be raised to ~1500°C which 
is 400-500°C higher than that of pure silica glass ware. Thermal shock 
resistance of these composites due to the surface compressive stress may 
be even higher than the pure silica glass. 



261 



Although the mechanical properties of these composites derived from 
chemical processes described in this dissertation are far from the 
theoretical limit, the possibility of closing the gap has been demon- 
strated to be real through the continuous understanding of the chemistry 
in the processes. The results are encouraging for the first attempts 
and values for af^g^^/p and Kj^/p are equal to or surpass all but the 
uniaxially oriented SiC fiber/glass hot-pressed composites. 

As the science of materials has developed it has moved more from 
giant leaps into the unknown to small steps forward in fairly clear di- 
rections. The small steps are many times more sophisticated and repre- 
sent a new level of scientific achievement. 

It is hoped that the small understandings of ceramics and molecular 
composites presented in this dissertation may accelerate the movements 
toward clearly defined targets which we know will be achieved. 

Some of more notable principles that were learned from this 
research for Chemically Derived Ceramic Composites may be summarized 
below. 

1) A precursor to SiC, known as polysilastyrene (PSS), had been 
shown to be potentially useful and superior to the existing precursor 
(polycarbosi 1 ane) if it can be crossl inked without adding oxygen to the 
polymer chain. First it was learned that PSS can be routinely synthe- 
sized in a one-step process with a relatively straight forward manner. 

2) Crossl inking of PSS can be achieved by using a chemical free 
radical initiator. Through the crossl inking process, production of a 
SiC monolithic body is possible without hot-pressing or high temperature 
sintering. 



262 



3) Increased crossl inkabil ity of an organosilane is achieved 
through reactive functional groups such as 

4) These organosil anes can be used to impregnate porous ceramic 
bodies such as silica gel followed by in situ crossl inking to yield a 
SiC/Si02 composite. The subsequent pyrolysis leads to hardening effect 
as well as a toughening and strengthening effect by the dispersed SiC in 
the surface layer of the matrix. 

5) The organosilane precursor can also be mixed with silica sol, 
followed by a cogellation, by an in situ crossl inking, and by a pyroly- 
sis to yield bulk SiC/SiiJ2 composites. 

6) There is(are) a chemical reaction(s) between the impregnated 
silane and the silica matrix which weaken the silica glass network. 
These chemical reaction(s) must be identified and controlled for improv- 
ed mechanical properties of silane/Si02 gel composites. 

7) The sol -gel process to obtain SiO^ glass can be used to form 
fibrous SiC/SiOo composites by dispersing the whiskers (or fibers) in 
the silica sol followed by castings, gellating, or cold pressing, and 
sintering. Characteristics of these composites are: a) low temperature 
processed (~200-5G0°C lower than the conventional hot-pressing process) 
so there is little damage on the fibers, and matrix and fiber inter- 
action is minimal; b) low density materials (~1. 7-2.1 g/cc); c) mediocre 
flexural strengths but good flexural strength to density ratio; d) sup- 
erior fracture toughness represented by the critical intensity factor, 
KjQ. Even more superior in Kj^ to density ratio to other ceramic com- 
posites of comparable composition, e) This remarkable Kjq value comes 



263 



from the load transfer from the matrix to the fiber as well as tne mech- 
anism of crack wandering and branching by pores. It was learned that 
the fracture toughness is limited by the diameter to length ratio of the 
fiber, f) Higher thermal shock resistance and better oxidation resist- 
ance than the SiC not in the composite, g) Complex shapes can be formed 
easily in plastic molds. The shapes and sizes are only limited by the 
molds. 

8) Molecular composite powders of Si02/SiC, Ti02(TiC)/SiC, and 
AlnOo/SiC can be made by mixing the two precursors of the respective 
phase with in situ crosslinking by a chemical free radical initiator. 
These composites appear to have a high devitrification temperature as 
well as high purity. 

9) These technques: crosslinking, impregnation, sol-gel processing 
for monolithic ceramics or glasses, and molecular level mixing of two or 
more precursors can be applied to many other ceramic composite systems. 

10) Silicon carbide incorporated into a porous silica glass matrix 
is more oxidation resistant than raw SiC. The SiC in the composite 
appears to be oxidized via an intermediate phase SiOC. 



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To be presented at the Engineering Ceramic Conf., Cocoa Beach, FL, 
Jan. 1986. 



BIOGRAPHICAL SKETCH 

Burtrand Insung Lee attended and yraduated from Montgomery Blair 
High School in Silver Spring, Maryland and Southern College in 
Collegedale, Tennessee, in 1971 and 1975, respectively. Upon graduation 
from Southern College as Chemistry major, he was employed oy Biospherics 
Inc., in Rockville, Maryland, as a chemist. At Biospherics he was in- 
directly engaged in the Viking Project search of life on Mars. 

After the decision to pursue full-time graduate work, he enrolled 
and graduated from Western Michigan University majoring in analytical 
chemistry in 1979. After some experiences in teaching and research in 
chemistry at State University of Mew York, he entered the University of 
Florida as a doctoral student. 



272 



I certify that I have read this study and that in my opinion it 
conforms to acceptable standards of scholarly presentation and is fully 
adequate in scope and quality, as a dissertation for the degree of 
Doctor of Philosophy. 




Larry L. Hench, Chairman 
Professor of Materials 
and Engineering 



Science 



I certify that I have read this study and that in my opinion it 
conforms to acceptable standards of scholarly presentation and is fully 
adequate in scope and quality, as a dissertation for the degree of 
Doctor of Philosophy. 




Christopher D. Batich 

Associate Professor of Materials 

Science and Engineering 



I certify that I have read this study and that in my opinion it 
conforms to acceptable standards of scholarly presentation and is fully 
adequate in scope and quality, as a dissertation for the degree of 
Doctor of Philosophy. 



Mr UhCL, 



Michael D. Sacks 

Associate Professor of Materials 

Science and Engineering 



I certify that I have read this study and that in my opinion it 
conforms to acceptable standards of scholarly presentation and is fully 
adequate in scope and quality, as a dissertation for the degree of 
Doctor of Philosophy. 



Lawrence E. Malvern 

Professor of Engineering Science 



I certify that I have read this study and that in my opinion it 
conforms to acceptable standards of scholarly presentation and is fully 
adequate in scope and quality, as a dissertation for the degree of 
Doctor of Philosophy. 




David E. Clark 

Associate Professor of Materials 

Science and Engineering 



This dissertation was submitted to the Graduate Faculty of the College 
of Engineering and to the Graduate School, and was accepted as partial 
fulfillment of the requirements for the degree of Doctor of Philosophy. 



May 1986 




Dean, Graduate School